Drawn heat treated steel wire for high strength spring use and pre-drawn steel wire for high strength spring use

ABSTRACT

Drawn heat treated steel wire for high strength spring use is provided containing, by mass %, C: 0.67% to less than 0.9%, Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%, and Al: 0.0005% to 0.003%, having Si and Cr satisfying the following formula: 
       0.3%≦Si−Cr≦1.2%,
 
     and having a balance of iron and unavoidable impurities, having as impurities, P: 0.025% or less and S: 0.025% or less, furthermore having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm, further having, as a metal structure, at least residual austenite in a volume rate of over 6% to 15%, having a prior austenite grain size number of #10 or more, and having a circle equivalent diameter of undissolved spherical carbides of less than 0.2 μm.

TECHNICAL FIELD

The present invention relates to drawn heat treated steel wire for highstrength spring use which can be used as a material for high strengthsprings produced by cold coiling and to pre-drawn steel wire.

BACKGROUND ART

The springs which are used for automobile engines, clutches, etc. arebeing required to offer more advanced performance and higher durabilityin order to deal with the trend toward lighter weights and higherperformances of automobiles. For this reason, their materials, that is,drawn heat treated steel wire for high strength spring use, are alsobeing required to offer high material strength. In general, whenproducing such small sized, high strength springs, the material of thedrawn heat treated steel wire for high strength spring use is quenchedand tempered to impart higher material strength in the drawn heattreated steel wire for high strength spring use, then is cold coiled toobtain a coil spring shape. Furthermore, stress-relief annealing orother heat treatment and nitriding are performed to obtain a finishedcoil spring. For this reason, drawn heat treated steel wire for highstrength spring use is required to have not only high strength, but alsoto have a high enough workability that it will not break at the coldcoiling and to suppress softening due to the annealing, nitriding, andother heat treatment performed after coiling, that is, to have tempersoftening resistance.

A spring is required to have fatigue characteristics, so drawn heattreated steel wire for high strength spring use is used as a materialand further nitrided or shot peened to raise the hardness of the surfacelayer of the spring. The durability of a spring includes fatiguecharacteristics and a sag property. The fatigue characteristics areaffected by the surface layer hardness. The sag property (property ofthe spring ending up plastically deforming in the load direction duringuse) is greatly affected by not only the surface layer hardness, butalso the hardness of the base material of the spring. For this reason,in steel wire for high strength spring use, the surface layer hardnessafter nitriding and the temper softening resistance at the inside wherenitrogen is not introduced by nitriding are important.

Furthermore, when producing a spring by cold coiling, when producing thematerial of the drawn heat treated steel wire for high strength springuse, oil tempering, induction hardening treatment, etc. where rapidheating and rapid cooling are possible may be used.

For this reason, the drawn heat treated steel wire for high strengthspring use can be reduced in prior austenite grain size, so a springwith excellent fracture characteristics can be obtained. However, ifdrawn heat treated steel wire for high strength spring use becomeshigher in strength, in cold coiling, breakage may occur and the springshape may not be able to be formed.

To deal with this problem, some of the inventors proposed drawn heattreated steel wire for high strength spring use obtained by controllingthe carbides, making the prior austenite finer, and achieving bothstrength and cold coiling ability (PLT 1). Furthermore, they proposeddrawn heat treated steel wire for high strength spring use obtained bycontrolling the residual austenite and carbides, refining the prioraustenite, and achieving both strength and cold coiling ability (PLT 2to PLT 4). In particular, the starting points of fracture caused by theformation of coarse oxides and carbides are suppressed and thedistribution of fine carbides of cementite required for securingstrength is made uniform so as to suppress deterioration of the fatiguecharacteristics and workability of the drawn heat treated steel wire forhigh strength spring use.

PLT 2 focuses on the fact that the region of sparse spherical carbideswith a circle equivalent diameter of 2 μm or more in the region of asparse distribution of fine spherical carbides (in particular,cementite) affects the dynamic characteristics and defines that region.

PLT 3 and PLT 4 take note of the effect of precipitation of finecarbides due to the addition of the alloy element V and limits thenitrogen (N) content to suppress undissolved spherical carbides. Thatis, they utilize the effect of precipitation of carbides, nitrides, andcarbonitrides of V to enable utilization for hardening the steel wire atthe tempering temperature or hardening the surface layer in nitriding.Furthermore, there is also an effect on suppressing coarsening of theaustenite grain size due to the formation of precipitates. The effect ofaddition of V is remarkable. However, undissolved carbides or nitrideeasily form, so even if suppressing the nitrogen (N), the control ofprecipitation has to be performed precisely.

Therefore, PLT 4 quantitatively compares the undissolved sphericalcarbides and the precipitated carbides and defines the amounts so as toobtain as much precipitated V carbides, which are effective for thefinal spring performance, as possible. Specifically, it proposes toweigh the residue of V carbides in the electrolytic solution at aconstant potential and compare this with the amount of V which passesthrough the filter (amount of precipitated V).

CITATION LIST Patent Literature

-   PLT 1: Japanese Patent Publication (A) No. 2002-180198-   PLT 2: Japanese Patent Publication (A) No. 2006-183137-   PLT 3: Japanese Patent Publication (A) No. 2006-342400-   PLT 4: International Publication WO2007/114491

SUMMARY OF INVENTION Technical Problem

In recent years, to raise the durability of high strength springs,surface hardening by nitriding has become a general practice.Furthermore, increasing the nitrided depth and shortening the nitridingtime by raising the treatment temperature is being studied. For thisreason, drawn heat treated steel wire for high strength spring use isbeing required to be further improved in temper softening resistance.

That is, a further better cold coiling ability than even withconventional drawn heat treated steel wire for high strength spring use,excellent temper softening resistance even after being held at 500° C.for 1 hour, internal softening kept to a minimum, and greater hardnessof the surfacemost layer are being sought.

The above conventional drawn heat treated steel wire for high strengthspring use secures a certain extent of uniform dispersion of finecarbides for improving the fatigue characteristics and workability.However, to improve the temper softening resistance, further uniformdispersion is necessary. In particular, the addition of V proposed inPLT 3 and PLT 4 does indeed have the effect of hardening the steel wireat the tempering temperature, hardening the surface layer in nitriding,and refining the austenite. However, on the other hand, control of thenitrogen (N) content is not easy. As a result, coarse carbides,nitrides, and carbonitrides are precipitated and cause degradation inthe fatigue strength.

PLT 3 adds Nb and Ti with the aim of the effect of trapping excessnitrogen (N). However, even if doing this, control to a suitable amountof N content is still not easy.

PLT 4 samples the residue of undissolved spherical carbides obtained asa result and compares it with the dissolved carbides. Therefore, it doesnot proactively control uniform dispersion of fine carbides.

Due to the above, the present invention has as its object to keep to aminimum the addition of V and other alloy elements, that is, withoutprecisely controlling the N content, develop drawn heat treated steelwire for high strength spring use which has excellent yield strength andhardness and excellent workability and which has superior surface layerhardness and internal hardness even after nitriding.

Further, as described in PLT 3 and PLT 4, to obtain excellent yieldstrength and hardness and excellent workability, the size of theundissolved spherical carbides in the steel should be small. Theeffective size is preferably 0.1 μm or less. If over 1 μm, thecontribution to strength and workability is lost and the deformationcharacteristics are just degraded. For this reason, the density ofpresence of undissolved spherical carbides with a circle equivalentdiameter of 0.2 μm or more becomes an important indicator. Therefore,the present invention has as its object the development of steel wirefor high strength spring use not allowing the presence of undissolvedspherical carbides with a circle equivalent diameter of 0.2 μm or more.

Solution to Problem

The inventors engaged in intensive research to solve the above problemsand as a result obtained the following discoveries:

(a) It was discovered that by strictly controlling the contents of C,Si, Mn, and Cr in the steel wire to suppress the formation of sphericalcarbides and by utilizing the residual austenite, even without addingalloy elements such as V, the drawn heat treated steel wire for highstrength spring use is improved in strength and cold coiling abilitycompared with the conventional.

(b) It was also discovered that by adding both Cr and Si in the steelwire in suitable amounts, the formation of undissolved sphericalcarbides and the softening in annealing or nitriding after coiling aresuppressed, and, furthermore, greater hardness of the nitrided layer canbe achieved.

That is, for increasing the strength in the fatigue characteristics,addition of Cr is effective, but Cr is an element which easily leavesbehind undissolved spherical carbides which would have a detrimentaleffect on the cold coiling ability. For this reason, the amount ofaddition had to be restricted. The inventors also took note of Si whichsuppresses the growth of undissolved spherical carbides and theformation of cementite. They discovered that if adding Si and togetherincreasing the amount of addition of Cr, the drawn heat treated steelwire can be increased in strength. Quantitatively, it is sufficient toadd large amounts of both Si and Cr and, as the relationship betweenthem, control the difference in amount of addition of Si and the amountof addition of Cr, that is, (Si−Cr) %.

(c) Further, it was discovered that by heating the bloom to 1250° C. ormore, it is possible to make Cr and other alloy elements in the steelmaterial uniformly disperse and suppress the formation of coarseundissolved spherical carbides and, furthermore, make fine carbidesuniformly disperse.

Undissolved spherical carbides are present in the steel material justafter casting and become causes of not only poor coiling ability, butalso breakage in rolling and drawing. For this reason, to prevent adetrimental effect in the steps of blooming after casting, wire rodrollingwire rod, patenting, quenching, and drawing, it is effective toraise the heating temperature in each step and constantly suppressundissolved spherical carbides.

(d) Furthermore, it was discovered that the addition of V has adetrimental effect on the mechanical properties and fatigue strength ofsteel wire for spring use.

That is, from just after casting to being worked into a spring, a steelmaterial is repeatedly heated. Usually, the undissolved sphericalcarbides are mainly cementite (Fe₃C). However, by repeating the heating,the undissolved spherical carbides often include Cr, V, etc. It islearned that not only are Cr, V, and other alloy elements wastefullyconsumed, but there is also a possibility of degrading the mechanicalcharacteristics after nitriding (surface hardness, internal hardness,etc.)

Further, as explained above, with the addition of V, control of thenitrogen (N) content is not easy. As a result, coarse carbides,nitrides, and carbonitrides precipitate and become causes of degradationin fatigue strength.

From these facts, the inventors discovered that by not adding V, or anextremely small amount even though added, and further, as explainedabove, controlling the amount of Cr in balance with the amount of Si, itis possible to suppress coarsening of the undissolved sphericalcarbides.

Here, “undissolved spherical carbides” means undissolved carbides with aratio of the maximum size (long size) and minimum size (short size)(aspect ratio) of 2 or less. Actually, “carbides” and “sphericalcarbides” are also undissolved. Here, in the sense of emphasis, whilesynonymous, these respectively are also called “undissolved carbides”and “undissolved spherical carbides”.

The present invention was made based on these discoveries. The gist ofthe invention is as follows:

(1) Pre-drawn steel wire for high strength spring use characterized bycontaining, by mass %,

C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%,and Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,having P and S as impurities comprisingP: 0.025% or less andS: 0.025% or less, and, furthermore,having a circle equivalent diameter of undissolved spherical carbides ofless than 0.2 μm.

(2) Pre-drawn steel wire for high strength spring use as set forth in(1) characterized by, further, containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%, W: 0.05 to 0.30%

Mg: 0.002% or less,Ca: 0.002% or less, andZr: 0.003% or less,when containing Vsatisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and, when containing Moand W,satisfying 0.05%≦Mo+W≦0.5%.(3) Drawn heat treated steel wire for high strength spring usecharacterized by containing, by mass %,C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%,and Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%, and

having a balance of iron and unavoidable impurities, having P and S asimpurities comprisingP: 0.025% or less andS: 0.025% or less,furthermore,having a metal structure comprised of at least residual austenite in avolume rate of over 6% to 15%,having prior austenite grain size number of #10 or more, andhaving a circle equivalent diameter of undissolved spherical carbides ofless than 0.2 μm.(4) Drawn heat treated steel wire for high strength spring use as setforth in (3) characterized by, further, containing, by mass %, one ormore of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%, W: 0.05 to 0.30%

Mg: 0.002% or less,Ca: 0.002% or less, andZr: 0.003% or less,when containing Vsatisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,when containing Mo and W,satisfying 0.05%≦Mo+W≦0.5%.(5) Drawn heat treated steel wire for high strength spring use as setforth in (3) or (4) characterized in that said drawn heat treated steelwire for high strength spring use has a tensile strength of 2100 to 2400MPa.(6) Drawn heat treated steel wire for high strength spring use as setforth in any one of (3) to (5) characterized in that said drawn heattreated steel wire for high strength spring use has a yield stress of1600 to 1980 MPa.(7) Drawn heat treated steel wire for high strength spring use as setforth in any one of (3) to (6) characterized said drawn heat treatedsteel wire for high strength spring use has a a surface Vicker'shardness of HV750 or more and an internal Vicker's hardness of HV570 ormore aftersoft nitriding of keeping at 500° C. for 1 hour.(8) A method of production of pre-drawn steel wire for high strengthspring use characterized by taking a bloom containing, by mass %,C: 0.67% or greater and less than 0.9%,

Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%,and Al: 0.0005% to 0.003%,

having Si and Cr satisfying the following formula:

0.3%≦Si−Cr≦1.2%,

having a balance of iron and unavoidable impurities,having P and S as impurities comprisingP: 0.025% or less andS: 0.025% or less, heating the bloom to 1250° C. or more, then hotrolling the bloom to produce a billet and heating the billet to 1200° C.or more, then hot rolling to produce pre-drawn steel wire.(9) A method of production of pre-drawn steel wire for high strengthspring use as set forth in (8) characterized by the bloom further,containing, by mass %, one or more of

V: 0.03 to 0.10%,

Nb: 0.015% or less

Mo: 0.05 to 0.30%, W: 0.05 to 0.30%

Mg: 0.002% or less,Ca: 0.002% or less, andZr: 0.003% or less,when containing Vsatisfying 1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and,when containing Mo and W,satisfying 0.05%≦Mo+W≦0.5%.(10) A method of production of pre-drawn steel wire for high strengthspring use characterized by further heating pre-drawn steel wire as setforth in (8) or (9) to 900° C. or more, then patenting it at 600° C. orless.(11) A method of production of heat treated steel wire for high strengthspring use characterized by drawing said pre-drawn steel wire which wasproduced by the method of production of pre-drawn steel wire as setforth in any one of (8) to (10), heating by a heating rate of 10° C./secor more up to an A₃ point, holding at a temperature of the A₃ point ormore for 1 minute to 5 minutes, then cooling by a cooling rate of 50°C./sec or more down to 100° C. or less.(12) A method of production of heat treated steel wire for high strengthspring use as set forth in (11) characterized by further holding andtempering it at 400 to 500° C. for 15 minutes or less.

Advantageous Effects of Invention

According to the present invention, in particular, due to the excellentcold coiling ability and temper softening resistance, even withnitriding at 500° C. for 1 hour, drawn heat treated steel wire for highstrength spring use with a high surface layer hardness and internalhardness and, furthermore, high strength spring excellent in durabilitycan be provided. The contribution in industry is extremely great.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is a micrograph of the metal structure showing one example ofspherical carbides in the drawn heat treated steel wire for highstrength spring use of the present invention. At the tips of the arrowsin the figure, undissolved spherical carbides are observed.

FIG. 2 is a view showing the shape of a punch for providing a notch in atest piece.

FIG. 3 is a view showing the step of providing a notch in a test piece.

FIG. 4 is a view showing an outline of a notch bending test.

FIG. 5 is a view showing a method of measurement of a notch bendingangle.

DESCRIPTION OF EMBODIMENTS

In general, a wire rod for a spring is produced as follows: Of course,production of springs is not limited to this here described process.This is described as just one example.

A bloom made of steel containing predetermined chemical compositions isrolled to obtain a billet. Next, the billet is rolled to produce apredetermined diameter of steel wire. The steel wire which is producedat this stage is called the “pre-drawn steel wire”.

The steel wire which is produced after rolling is patented and drawn toobtain further finer steel wire, then the working stress at the surfacelayer is removed and subsequent cold coiling workability is obtained byheat treatment (quenching and tempering). The steel wire which isproduced at this stage is called the “drawn heat treated steel wire”.

Next, the spring is worked by cold coiling and is improved in strengthand surface hardness by nitriding. In this way, a “spring” is producedas a final product.

First, the chemical compositions of the drawn heat treated steel wirefor high strength spring use of the present invention and its material,that is, pre-drawn steel wire for high strength spring use, will beexplained. Here, the “%” in the chemical compositions means mass %except when otherwise indicated.

C: 0.67% to less than 0.9%

C is an important element which has a great effect on the strength ofthe steel material and contributes to the formation of residualaustenite as well. In the present invention, to obtain sufficientstrength, the lower limit of the amount of C is made 0.67% or more. Toraise the strength, the amount of C is preferably made 0.70% or more,more preferably 0.75% or more.

On the other hand, if the amount of C becomes 0.9% or more, excessivecoprecipitation results, a large amount of coarse cementite isprecipitated, and the toughness remarkably falls. Further, if the amountof C is excessive, coarse spherical carbides are formed and the coilingability is impaired. Therefore, the upper limit of the amount of C ismade less than 0.9%. From the viewpoint of suppressing the formation ofspherical carbides, the upper limit of the amount of C is preferably0.85%, more preferably 0.80%.

Si: 2.0 to 3.5%

Si is an important element for improving the temper softening resistanceof the steel and the sag property of the spring. To obtain theseeffects, 2.0% or more has to be added. Further, Si is effective forspheroidization and refinement of the cementite. To suppress theformation of coarse spherical carbides, 2.1% or more of Si is preferablyadded. To raise the internal hardness after nitriding and othertreatment for making the surface layer harder, 2.2% or more of Si ismore preferably added. Furthermore, from the balance with Cr, Si is morepreferably made 2.3% or more. Si is sometimes made 3.0% or more.

On the other hand, if excessively adding Si, the steel wire hardens andbecomes brittle, so the upper limit of the amount of Si is made 3.5% orless. From the viewpoint of the prevention of embrittlement, the upperlimit is preferably made 3.4%, more preferably 3.3% or less.

Mn: 0.5 to 1.2%

Mn is an element which is important for raising the quenchability andstably securing the amount of residual austenite. In the presentinvention, to raise the yield strength of the steel wire and secure theresidual austenite, Mn has to be added in 0.5% or more, more preferably0.65% or more, still more preferably 0.70% or more.

On the other hand, if excessively adding Mn, the residual austeniteincreases. In working, work-induced martensite is formed and the coldcoiling ability is impaired. To prevent embrittlement due to excessiveaddition of Mn, the upper limit of the amount of Mn is made 1.2% orless, preferably 1.1% or less, more preferably 1.0% or less.

Cr: 1.3 to 2.5%

Cr is an element which is effective for improving the quenchability andtemper softening resistance. To obtain these effects, 1.3% or more of Crhas to be added. When performing the nitriding, it is possible to makethe hardened layer obtained by nitriding deeper by the addition of Cr.Therefore, when imparting hardening by nitriding and softeningresistance at the nitriding temperature, over 1.5% of Cr is preferablyadded. More preferably, 1.7% or more should be added.

On the other hand, if the amount of Cr is excessive, the production costbecomes higher. Not only this, dissolution of the carbides is impaired,undissolved spherical carbides become greater, and the coiling abilityis impaired, so the upper limit of the amount of Cr is made 2.5% orless. Further, if the amount of Cr is large, to suppress formation ofcoarse cementites, the amount of Cr is preferably suppressed to 2% orless. Furthermore, to obtain both strength and coiling ability, theupper limit of the amount of Cr is preferably made 1.8% or less.

N: 0.003 to 0.007%

N is an element, in the present invention, which forms nitrides with Aletc. included as impurities in the steel. To utilize the fine nitridesand refine the prior austenite, 0.003% or more of N has to be included.On the other hand, if the amount of N is excessive, the nitrides coarsenand the cold coiling ability and fatigue characteristics fall.Therefore, the upper limit of the amount of N is made 0.007% or less.Further, if considering the ease of heat treatment etc., the amount of Nis preferably 0.005% or less.

P: 0.025% or less

P is an impurity. It causes the steel to harden, forms segregation, andcauses embrittlement, so the upper limit of the amount of P is made0.025% or less. Further, the P which segregates at the prior austenitegrain boundaries causes the toughness and delayed fracture resistanceetc. to fall, so the upper limit of the amount of P is preferably made0.015% or less. Furthermore, when the yield strength of the steel wirewill exceed 2150 MPa, the amount of P is preferably limited to less than0.010%.

S: 0.025% or less

S is also an impurity. If present in steel, it causes the steel toembrittle, so the upper limit of the amount of S is made 0.025% or less.To suppress the effect of S, addition of Mn is effective. However, MnSis an inclusion. In particular in high strength steel, MnS sometimesbecomes starting points of fracture. Therefore, to suppress theoccurrence of fracture, the upper limit of the amount of S is preferablymade 0.015% or less. Furthermore, when the yield strength of the drawnheat treated steel wire for high strength spring use will exceed 2150MPa, the amount of S is preferably limited to less than 0.01%.

Al: 0.0005 to 0.003%

Al is a deoxidizing element. It affects the formation of oxides. Ifforming hard oxides, the fatigue durability falls. In particular, inhigh strength springs, if excessively adding Al, the fatigue strengthfluctuates and the stability is impaired. In the drawn heat treatedsteel wire for high strength spring use of the present invention, if theamount of Al exceeds 0.003%, the rate of occurrence of fracture due toinclusions becomes greater, so the amount of Al is limited to 0.003% orless. The upper limit value of the amount of Al is preferably 0.0028%,more preferably 0.0025%.

On the other hand, if the amount of Al becomes less than 0.0005%,silica-based hard oxides are easily formed. For this reason, the amountof Al is made 0.0005% or more. The lower limit of the amount of Al ispreferably 0.0007%, more preferably 0.0008%, further preferably 0.001%or more.

Next, the point of the present invention, that is, the relationshipbetween Si and Cr, will be explained. It is already known that Si and Crare both important for increasing the strength of spring steel. However,excessive addition causes problems.

0.3%≦Si−Cr≦1.2%

If the amount of Si exceeds the prescribed amount, the embrittlementbecomes extreme and the workability in coiling is impaired. Not onlythat, decarburization in the intermediate processes becomes remarkable.For this reason, in the final product of the spring, the surface layerhardness becomes lower and the durability falls. Further, decarburizedparts are randomly formed, so the stability of the strength of thespring product is impaired. When the amount of Si is smaller than theprescribed amount, the strength falls. Furthermore, the sag property isinsufficient. This appears in the hardness after nitriding as well.Sufficient hardness cannot be secured both at the surface layer andinside.

However, the relationship between Si and Cr through the cementite in thesteel is important. That is, Si is an element which destabilizescementite. If adding a large amount of Cr or other element whichstabilizes cementite, in heating, there is the effect of promoting theformation of a solid solution by the cementite. Therefore, regardless ofadding a large amount of Cr, if the amount of addition of Si is small,the amount of undissolved spherical carbides becomes greater and theworkability is remarkably reduced. The inventors discovered that it ispossible to use the difference between the Si content (mass %) and Crcontent (mass %) in the steel, that is, the Si−Cr amount, as ayardstick. That is, when the value of Si−Cr is smaller than 0.3%, theamount of Cr becomes relatively large and undissolved spherical carbideseasily remain. On the other hand, if over 1.2%, Si becomes relativelyexcessive and easily causes embrittlement, decarburization, or otherproblems. Therefore, the value of Si−Cr should be made 0.3 to 1.2%.

From the viewpoint of suppressing the formation of carbides, a largeramount of Si−Cr enables the undissolved carbides to be suppressed, butindustrially, if the Si is too great, the depth of the hardened layerformed by the nitriding will easily become shallow. For this reason, ifconsidering the behavior of undissolved spherical carbides and thehardened layer formed by nitriding, preferably Si−Cr≦0.9%, morepreferably Si—Cr≦0.75%. Further, from the viewpoint of relativelyreducing the amount of Cr and reducing the residual presence ofundissolved spherical carbides, the lower limit is preferably0.35≦Si−Cr, more preferably 0.4≦Si−Cr.

Next, the selectively added chemical compositions will be explained.

V: 0.03 to 0.10%

V is an element which forms nitrides, carbides, and carbonitrides. FineV nitrides, carbides, and carbonitrides with a circle equivalentdiameter of less than 0.2 μm are effective for refinement of the prioraustenite. Further, they may also be utilized for hardening the surfacelayer by nitriding. However, on the other hand, undissolved carbides andnitrides are easily formed, so even if suppressing the nitrogen (N), itis necessary to precisely control the precipitation.

For this reason, in the present invention, V is not deliberately added.

To obtain such an effect of addition of V, a fine amount can be added.To obtain these effects, V should be added in 0.03% or more, preferably0.035% or more, more preferably 0.04% or more.

On the other hand, if adding over 0.10% of V, coarse spherical carbidesare formed and the cold coiling ability and spring fatiguecharacteristics are impaired. Therefore, the V content should be made0.1% or less. Further, by the addition of V, before drawing, asupercooled structure causing cracks and breakage in drawing easily isformed. For this reason, the upper limit of the amount of V ispreferably made 0.09% or less, more preferably 0.08% or less, mostpreferably 0.05% or less. In particular, in the case of adding a fineamount of Nb, the amount of addition of V is preferably made 0.05% orless. Further, V is an element which greatly affects the formation ofresidual austenite in the same way as Mn, so the amount of V has to beprecisely controlled together with the amount of Mn.

Nb: 0.015% or less

Nb is an element which forms nitrides, carbides, and carbonitrides insteel. These precipitates are sometimes used for control of theaustenite grain size etc. However, simultaneously, excessive additionreduces the ductility when hot and results in easier cracking in rollingor hot forging. For this reason, excessive addition must be avoided.

Nb is added for the purpose of controlling the amount of N. Theprecipitates are not directly used for controlling the quality. Valvesprings and other springs are produced by quenching, tempering, thencold coiling, but at that time, the dissolved nitrogen obstructs colordeformation and reduces the limit strain. For this reason, the coilingability is impaired. Therefore, by adding Nb and forming nitrides at ahigh temperature, there is the effect that the dissolved nitrogen in thesteel in the steel matrix is lowered and the cold workability isimproved.

Further, the addition of a fine amount of Nb is also effective forsuppressing V and other undissolved spherical carbides mixed in asunavoidable impurities. V is an element which is effective for improvingthe temper softening resistance in nitriding and the surfacemost layerhardness. However, if the amount of addition becomes greater, even inthe patenting, quenching, and other heating for obtaining an austenitephase for producing drawn heat treated steel wire for high strengthspring use, V nitrides, V carbides, and V carbonitrides often are notsufficiently dissolved. The undissolved spherical carbides of V growfrom cores of the V-based nitrides formed at the time of normal hightemperature. As a result, undissolved spherical carbides remain and thecoiling ability is impaired. For this reason, when suppressing theundissolved spherical carbides, it is necessary to suppress the amountof addition of V. In the present invention, V was not made an essentialelement.

As opposed to this, Nb forms nitrides at a higher temperature than V.For this reason, in the steelmaking process, addition of Nb suppressesthe formation of V nitrides. That is, Nb forms nitrides in the hightemperature region where V dissolves and does not form nitride.Furthermore, at the high temperature where V nitrides are formed, Nbconsumes nitrogen, so V nitrides become harder to form even when cooled.For this reason, the addition of a fine amount of Nb is particularlyeffective for suppressing undissolved spherical carbides and securingcoiling ability when adding a large amount of V.

If the amount of addition of Nb is over 0.015%, the hot ductility isimpaired and the occurrence of defects and other problems in rollingbecomes easy. For this reason, the amount of addition is made 0.015% orless, preferably 0.010% or less, more preferably 0.005% or less, mostpreferably less than 0.001%.

On the other hand, the effect of Nb in controlling the amount of N inspring steel appears starting from 0.0005%, so when adding Nb, 0.0005%or more is preferably added. Further, when adding V etc., addition of afine amount of Nb is more effective. A range of 0.003 to 0.012% ispreferable. Furthermore, a range of 0.005 to 0.009% is more preferable.The effect is obtained even at 0.005 to 0.001%.

1.4%≦Cr+V≦2.6%

In the present invention, V is not deliberately added. However, asexplained above, addition of a fine amount of V has an effect on therefinement of the prior austenite and formation of residual austenite.By precisely controlling the sum of the amounts of addition of Cr and Vwith respect to V, it is possible to raise the strength to make thesurface layer hardness after nitriding and the internal hardnesssuitable for high strength springs.

Cr and V are both elements which prevent softening upon the heating bythe annealing or nitriding etc. performed after the spring coiling, thatis, impart so-called temper softening resistance. In particular,nitriding causes nitrides to precipitate at the nitrided part of thesurface layer to thereby improve the surface hardness and increase thenitriding effect. Further, even at the inside where nitriding does notspread, decomposition of the carbides is suppressed. Further, there isthe effect of suppressing softening by precipitation of carbides. On theother hand, both are elements which facilitate the formation ofundissolved spherical carbides. Cr dissolves in the cementite toincrease the stability so in the heating steps for dissolving thecementite (heating at time of patenting and heating at time ofquenching), suppresses the dissolution of the cementite, often remainsas undissolved spherical carbides. Further, V also has a dissolutiontemperature of the precipitates higher than the A3 point of steel, soeasily remains as undissolved spherical carbides.

If the total of the contents of Cr and V, that is, Cr+V, is less than1.4%, the surface layer hardness of the high strength spring falls belowHV750 and the internal hardness falls below HV570. For this reason, Cr+Vis preferably 1.4% or more. Furthermore, 1.5% or more is preferable. Onthe other hand, excessive addition of Cr+V of over 2.6% leaves behindlarge amounts of undissolved spherical carbides, so the coiling abilityis impaired. Therefore, 2.6% is made the upper limit. Further, the Cr+Vis preferably 2% or less, more preferably 1.8% or less.

0.7%≦Mn+V≦1.3%

Mn and V are elements which improve the quenchability and also have alarge effect on the formation of residual austenite. If Mn is largerthan the prescribed amount, a large amount of residual austeniteremains. Therefore, the sum of both Mn and the V which is included as anunavoidable impurity has a direct effect on the austenite behavior. Ifthese exceed their prescribed amounts, the amount of residual austeniteincreases. Not only is the workability affected, but also the yieldstrength is greatly affected. Sufficiently durability cannot be secured.

For this reason, in the present invention, the total of the contents ofMn and V, that is, Mn+V, is made 0.7 to 1.3%. To secure a volume rate ofover 6% of residual austenite, the lower limit of Mn+V has to be made0.7% or more.

As a result, transformation induced plasticity causes the ductility tobe improved and enables the cold coiling ability to be secured. On theother hand, to make the residual austenite a volume rate of 15% or less,the upper limit value of Mn+V has to be made 1.3% or less. Due to this,the formation of work-induced martensite due to strike marks in coldcoiling is suppressed and local embrittlement can be prevented.

Mo: 0.05 to 0.30%

Mo is an element which improves the quenchability. Further, it is alsoextremely effective for improving the temper softening resistance. Inthe present invention, in particular, to further improve the tempersoftening resistance, 0.05% or more of Mo can be added. Further, Mo isan element which forms Mo-based carbides in the steel. The temperatureat which the Mo-based carbides precipitate is lower than carbides of Vetc. For this reason, addition of a suitable amount of No is alsoeffective for suppressing coarsening of carbides. Addition of 0.10% ormore of Mo is preferable. On the other hand, if the amount of additionof Mo is over 0.30%, a supercooled structure easily forms in hotrolling, the patenting before drawing, etc. Therefore, to suppress theformation of a supercooled structure causing cracking or wire breakagein drawing, the upper limit of the amount of Mo is made 0.30% or less,preferably 0.25% or less. Further, if the amount of Mo is large, in thepatenting, the time until the end of the pearlite transformation becomeslonger, so the amount of Mo is preferably made 0.20% or less.Furthermore, to shorten the patenting time and stably end the pearlitetransformation, 0.15% or less is preferable.

W: 0.05 to 0.30%

W, like Mo, is an element which is effective for improvement of thequenchability and temper softening resistance and is an element whichprecipitates in the steel as carbides. In the present invention, inparticular, to improve the temper softening resistance, 0.05% or more ofW is added.

On the other hand, if excessively adding W, a supercooled structure isformed which causes cracking or wire breakage in drawing, so the amountof W has to be made 0.30% or less.

Furthermore, if considering the ease of heat treatment etc., the amountof W is preferably 0.1 to 0.2%, more preferably 0.13 to 0.18%.

0.05%≦Mo+W≦0.5%

Mo and W are elements which are effective for improvement of the tempersoftening resistance. If adding the two combined, the growth of carbidesis suppressed and the temper softening resistance can be remarkablyimproved compared with addition of Mo and W alone. In particular, toimprove the temper softening resistance in heating to 500° C., Mo+W hasto be made 0.05% or more, preferably 0.15% or more.

On the other hand, if Mo+W is over 0.5%, in hot rolling, patentingbefore drawing, etc., a so-called supercooled structure of martensite,bainite, etc. is formed. To suppress the formation of a supercooledstructure causing cracks or wire breakage in drawing, the upper limit ofMo+W is made 0.5% or less, preferably 0.35% or less.

Next, Mg, Ca, and Zr will be explained.

Mg: 0.002% or less

Mg forms oxides in molten steel higher in temperature than the MnSforming temperature. At the time of formation of MnS, it is alreadypresent in the molten steel. Therefore, it can be used as a nuclei forprecipitation of MnS. Due to this, the distribution of the MnS can becontrolled. Further, in number distribution as well, Mg-based oxides aremore finely dispersed in the molten steel compared with the Si- andAl-based oxides which are often seen in conventional steel, so MnSformed around cores of Mg-based oxides are finely dispersed in thesteel. Therefore, even with the same S content, depending on thepresence of Mg, the MnS distribution differs. Adding these makes the MnSgrain size finer. By making the MnS finely disperse, it is possible torender the action as a starting point of fatigue of MnS harmless. Theeffect is sufficiently obtained even in fine amounts. Preferably Mg0.0002% or more, more preferably 0.0005% or more, should be added.

However, with addition of over 0.001%, it is difficult for the Mg toremain in the molten steel, there is an effect on the oxide composition,and the rate of appearance of oxides as initiation sites of fatiguebecomes higher, so 0.002% is made the upper limit. Therefore, the upperlimit of the amount of addition of Mg was made 0.002%, preferably0.0015% or less. Furthermore, in the case of spring steel, compared withother steel for structural use, the amount of addition of S issuppressed, so if considering the yield etc., 0.001% or less ispreferable. Further, when used for a high strength valve spring, theinclusion susceptibility is high, so Mg has the effect of improving thecorrosion resistance and resistance to delayed fracture preventingrolling cracks due to the effect of the distribution of MnS etc.Addition of as much as possible is preferable, so control of the amountof addition in the extremely narrow range of 0.0002 to 0.001% ispreferable.

Ca: 0.002% or less

Ca is an oxide- and sulfide-forming element. In spring steel, it makesthe MnS spherical to thereby suppress the length of MnS serving asinitiation sites of fatigue and other fracture and render it harmless.The effect is similar to Mg. Addition of 0.0002% or more is preferable.Further, even if over 0.002% is added, not only is the yield poor, butalso oxides and CaS and other sulfides are formed and trouble inproduction and degradation in spring fatigue durability characteristicsare caused, so the amount was made 0.002% or less. Regarding the amountof addition, when used for a high strength valve spring, the inclusionsusceptibility is high, so the amount is preferably 0.0015% or less,more preferably 0.001% or less.

Zr: 0.003% or less

Zr is an oxide-, sulfide-, and nitride-forming element. In spring steel,the oxides are finely dispersed, so like with Mg, form nuclei forprecipitation of MnS and can make the MnS finely disperse. Due to this,it is possible to improve the fatigue durability and, further, increasethe ductility to thereby improve the coiling ability. 0.0002% or more ispreferably added. Further, even if over 0.003% is added, not only is theyield poor, but oxides and ZrN, ZrS, and other nitrides and sulfides areformed and trouble in production or degradation in the spring fatiguedurability characteristics is caused, so the amount is made 0.003% orless. The amount of addition is preferably 0.0025% or less. Furthermore,when used for high strength valve spring, there is also the effect thatthe coiling ability is improved by the control of sulfides, so additionis preferred, but to minimize the effects on the dimensions ofinclusions, suppression to 0.0015% or less is preferable.

Note that, the above optionally added chemical compositions, ifcontained in fine amounts, do not impair the effects of the steel wirecomprised of the basic chemical compositions of the present invention.

Next, the metal structure of the steel wire for high strength spring useof the present invention will be explained.

Undissolved Spherical Carbides

Undissolved spherical carbides perform the important role of securingstrength in steel wire for high strength spring use. On the other hand,the presence of undissolved spherical carbides causes the coilingability to deteriorate. Further, coarse carbides cause the fatiguecharacteristics to degrade as well. Therefore, suppressing undissolvedspherical carbides in coiling and causing uniform dispersion of finecarbides after the final nitriding are essential for solving the problemof the present invention.

The steel wire for high strength spring use of the present invention hasa long size of the undissolved spherical carbides of 0.2 μm or less,that is, is suppressed in coarsening. The undissolved spherical carbidesare already present after wire rod rolling (that is, the pre-drawn steelwire).

The undissolved spherical carbides are hard to go into solid-solution inthe subsequent heat treatment (patenting, generation of working heat indrawing, and quenching and tempering, for instance). Rather, theysometimes grow in these heat treatment steps and coarsen. That is, theundissolved spherical carbides in the pre-drawn steel wire sometimes actas nuclei for coarsening of themselves.

For this reason, to restrict the coarsened undissolved sphericalcarbides of the steel wire after heat treatment (heat treated steelwire), it is important to reduce as much as possible the undissolvedspherical carbides which are present in the pre-drawn steel wire. Due tothe above, the definition regarding the “undissolved spherical carbides”has important meaning in not only the pre-drawn steel wire for highstrength spring use according to the present invention, but also thedrawn heat treated steel wire for high strength spring use.

The steel wire for high strength spring use of the present invention isincreased in strength by having C added, having Mn and Cr added and,further having Mo, W, and other so-called alloy elements added. Whenadding large amounts of C and, in particular, Cr and other alloyelements which form nitrides, carbides, and carbonitrides, sphericalcementite carbides and alloy-based carbides easily remain in the steel.Spherical cementite carbides and alloy-based carbides are undissolvedspherical carbides which do not dissolve in the steel in heating in thehot rolling.

Note that, in the present invention, spherical alloy-based carbides andspherical cementite carbides will be referred to all together asspherical carbides. In the steel, there are pin-shaped carbidescorresponding to the pin-shaped structure of tempered martensite, butthese pin-shaped carbides are not included in the spherical carbides ofthe present invention. The pin-shaped carbides are not present rightafter quenching and precipitate in the process of tempering. Thetempered martensite structure is a structure suitable for achieving bothstrength and toughness and workability. Being pin-shaped is, in acertain sense, the ideal form in carbides.

Strictly speaking, if carbides with an aspect ratio of 2 or more(pin-shaped carbides) also coarsen, the workability may be impaired.However, in actuality, pin-shaped carbides become coarse when thetempering temperature is high or the holding time in tempering isextremely long. The effect on performance is to make the strength andhardness insufficient. Problems arise in different areas than withundissolved spherical carbides. In the 2100 MPa or so strength steelwire covered by the present invention, coarse pin-shaped carbides arenot formed. Therefore, in the present invention, pin-shaped carbides arenot covered. As explained above, the normally precipitated carbides areundissolved, but in the present invention, the term “undissolved” addedto the top. This just stresses the undissolved nature. In the presentinvention, “undissolved spherical carbides” and “spherical carbides” aresynonymous.

The undissolved spherical carbides can be observed under a scanningelectron microscope (SEM) by polishing a sample obtained from pre-drawnsteel wire or drawn heat treated steel wire for high strength spring useto a mirror finish and etching it by picral or electrolytically etchingit. Further, they can be observed by the replica method under atransmission type electron microscope (TEM).

FIG. 1 shows an example of a structural photograph of a sample afterelectrolytic etching as observed under an SEM. In the structuralphotograph of FIG. 1, the steel is observed to have two types ofstructures of the matrix, that is, pin-shaped structures and sphericalstructures. Among these, the pin-shaped structures are temperedmartensite formed by quenching and tempering. On the other hand, thespherical structures are carbides 1 made spherical by not dissolvinginto the steel due to the heating of the hot rolling and by being madespherical by quenching and tempering by oil tempering or inductionhardening treatment (undissolved spherical carbides). Spherical carbidescan be observed at the front end of the arrow in FIG. 1.

Circle Equivalent Diameter of Undissolved Spherical Carbides of LessThan 0.2 μm

In the present invention, the undissolved spherical carbides affect thecharacteristics of the drawn heat treated steel wire for high strengthspring use, so are controlled in size as follows: Note that, in thepresent invention, compared with the prior art, further finer sphericalcarbides are defined for achieving both higher performance andworkability. Spherical carbides with a circle equivalent diameter ofless than 0.2 μm are extremely effective for securing the strength andtemper softening resistance of the steel.

On the other hand, spherical carbides with a circle equivalent diameterof 0.2 μm or more do not contribute to improvement of the strength ortemper softening resistance and degrade the cold coiling ability. Forthis reason, the present invention is characterized by not allowing theformation of spherical carbides with a circle equivalent diameter of 0.2μm or more.

The pre-drawn steel wire and drawn heat treated steel wire of thepresent invention is characterized in that the undissolved sphericalcarbides have a circle equivalent diameter of less than 0.2 μm. For thisreason, it is possible to secure strength while securing workability aswell.

As explained above, the pre-drawn steel wire has to be then patented,drawn and heated, quenched and tempered, or otherwise heat treated, sothe undissolved spherical carbides may grow and coarsen. For thisreason, the circle equivalent diameter of the undissolved sphericalcarbides in the pre-drawn steel wire is preferably made smaller than 0.2μm.

From the results of experiments of the inventors, the circle equivalentdiameter of undissolved carbides of the pre-drawn steel wire isconfirmed to be able to be reduced to 0.18 μm or less. Further, it isalso confirmed that if making the billet heating temperature 1250° C. ormore, the diameter can be made 0.15 μm or less.

Here, the methods of measuring the circle equivalent diameter anddensity of presence of spherical carbides will be explained. A samplewhich is taken from steel wire for high strength spring use is polishedand electrolytically etched. Note that, the observed location israndomly selected near the center of the radius of the heat treated wirerod (steel wire), that is, the so-called “½R part”, so as to enableelimination of special conditions such as decarburization and centersegregation. Further, the measurement area is 300 μm² or more. Inelectrolytic etching, the surface of the sample is corroded byelectrolytic action in an electrolytic solution (a mixture of acetylacetone 10 mass %, tetramethyl ammonium chloride 1 mass %, and a balanceof methyl alcohol) using the sample as the anode and platinum as thecathode using a current generator with a lower potential. The potentialbecomes constant at a potential suitable for the sample in the range of−50 to −200 mV vs SCE. For the steel wire of the present invention, itis preferable that it become constant at −100 mV vs SCE.

The amount of power run can be found by the total surface area of thesample×0.133 [c/cm²]. Note that, when embedding the sample in a resin,not only the polished surface, but also the area of the sample surfacein the resin are added to the total surface area of the sample. Thepower starts to be run, then the sample is held for 10 seconds, then thepower is stopped and the sample is cleaned.

After that, the sample is observed under an SEM and a structuralphotograph of the spherical carbides is taken. Under the SEM, thestructures which appear relatively white and which have a ratio (aspectratio) of the maximum size (long size) and minimum size (short size) of2 or less are the spherical carbides. The magnification of thephotograph taken under the SEM is X1000 or more, while X5000 to X20000is preferable. For the measurement locations, 10 fields were randomlyselected from locations at a depth of about 0.5 to 1 mm from the surfaceof the wire rod while avoiding the center segregation parts. The thuscaptured SEM structural photographs were processed by image processingto measure the minimum size (short size) and maximum size (long size) ofthe spherical carbides seen in the measured fields and calculate thecircle equivalent diameter. The circle equivalent diameter is thediameter when calculating the area of an undissolved carbide in a fieldby image processing and converting it to a circle of the same area.Further, it is also possible to measure the density of presence ofspherical carbides with a circle equivalent diameter of 0.2 μm or moreseen in the measurement field.

Metal Structure of Pre-Drawn Steel Wire For High Strength Spring Use andDrawn Heat Treated Steel Wire

The metal structure of the drawn heat treated steel wire for highstrength spring use according to the present invention is comprised of,by volume rate, over 6% to 15% of residual austenite and a balance oftempered martensite. Fine inclusions are allowed. The “fine inclusion”are oxides and sulfides. The oxides are deoxidation products of Al andSi etc., while the sulfides correspond to MnS, CaS, etc. Further, thebalance of the tempered martensite structure also includes undissolvedspherical carbides in fine amounts.

The prior austenite grain size number in the structure is #10 or more,while the circle equivalent diameter of the spherical carbides is lessthan 0.2 μm.

Further, in the metal structure of the pre-drawn steel wire for highstrength spring use according to the present invention, the pearlitestructure accounts for 90% or more, preferably 95% or more, morepreferably 98% or more. A substantially 100% pearlite structure isideal.

Prior Austenite Grain Size Number: #10 or more

The drawn heat treated steel wire for high strength spring use of thepresent invention is mainly comprised of tempered martensite instructure. The prior austenite grain size has a great effect on thecharacteristics. That is, if refining the grain size of the prioraustenite, due to the effect of grain refinement, the fatiguecharacteristics and the coiling ability are improved. In the presentinvention, to obtain sufficient fatigue characteristics and coilingability, the prior austenite grain size number is made #10.

Refining the prior austenite is particularly effective for improving thecharacteristics of the drawn heat treated steel wire for high strengthspring use. The prior austenite grain size number is preferably made#11, more preferably #12. To refine the grain size of prior austenite,it is effective to lower the heating temperature of the quenching. Notethat, the “prior austenite grain size number” is based on JIS G 0551. Ifactually performing the quenching by lowering the heating temperatureand shortening the time, the prior austenite grain size can be refined,but unreasonable low temperature, short time treatment not onlyincreases the undissolved spherical carbides, but also sometimes resultsin insufficient austenite transformation itself and two-phase quenching.Conversely, sometimes the coiling ability and the fatiguecharacteristics are lowered. For this reason usually #13.5 is the upperlimit.

Residual Austenite: Over 6% to 15% (volume rate)

The microstructure at the drawn heat treated steel wire for highstrength spring use after quenching and tempering is comprised oftempered martensite, residual austenite, and a slight volume fraction ofinclusions (here, precipitates also expressed included in inclusions).Residual austenite is effective for improving the cold coiling ability.In the present invention, to secure the cold coiling ability, the volumerate of the residual austenite is made over 6%, preferably 7% or more,more preferably 8% or more.

On the other hand, if the residual austenite exceeds a volume rate of15%, the martensite which is formed due to the work-inducedtransformation causes the cold coiling characteristics to drop.Therefore, the volume rate of the residual austenite is made 15% orless, preferably 14% or less, more preferably 12% or less.

The volume rate of the residual austenite can be found by the X-raydiffraction method and the magnetic measurement method. Among these, themagnetic measurement method enables simple measurement of the volumerate of the residual austenite, so is the preferable measurement method.Here, the volume rate is measured, but the obtained figures are the sameas the area rate.

Note that, residual austenite is softer than tempered martensite, soreduces the yield strength. Further, the transformation inducedplasticity is used to improve the ductility, so this remarkablycontributes to improvement of the cold coiling ability. On the otherhand, residual austenite often remains at the segregated parts, prioraustenite grain boundaries, and near regions sandwiched by the saggrains, so the martensite which is formed by the work-inducedtransformation (work-induced martensite) becomes starting points offracture. Further, if the residual austenite increases, the temperedmartensite falls relatively.

For this reason, in the past, the drop in the strength and cold coilingability due to the residual austenite had been considered an issue.However, in high strength steel wire of over 2000 MPa, the amounts ofaddition of C, Si, Mn, Cr, etc. become greater, so for improvement ofthe cold coiling ability, utilization of the residual austenite isextremely effective. Further, recently, high precision spring workingtechnology has made it possible to suppress the deterioration of thecoiling characteristics even if high hardness parts are locally formeddue to the work-induced martensite formed in shaping the spring.

Next, the mechanical properties of the drawn heat treated steel wire forhigh strength spring use of the present invention will be explained.

To reduce the size and lighten the weight of a spring, it is effectiveto make it higher in strength. Further, a spring is required to have asuperior fatigue strength. In the present invention, a high strengthspring is produced by bending the material of the drawn heat treatedsteel wire for high strength spring use to a desired shape, thennitriding, shot peening, or otherwise hardening the surface. In thenitriding, the spring is heated to 500° C. or so, so the spring issometimes softened more than the material of the drawn heat treatedsteel wire for high strength spring use.

Therefore, to raise the strength of the spring and improve the fatiguecharacteristics, it is necessary to secure the yield strength of thematerial of the drawn heat treated steel wire for high strength springuse. Further, in order for the drawn heat treated steel wire for highstrength spring use to be worked into the desired shape of a spring,cold coiling ability is demanded, so the upper limit of the yieldstrength has to be limited.

Yield Strength: 2100 to 2400 MPa

If the drawn heat treated steel wire for high strength spring use ishigh in yield strength, it is possible to improve the fatiguecharacteristics and sag property of the spring hardened at the surfaceby nitriding etc. In the present invention, to improve the fatiguecharacteristics and sag property of the spring, the yield strength ofthe drawn heat treated steel wire for high strength spring use is made2100 MPa or more.

Further, the higher the drawn heat treated steel wire for high strengthspring use in yield strength, the better the spring in fatiguecharacteristics, so the drawn heat treated steel wire for high strengthspring use has a yield strength of preferably 2200 MPa or more, morepreferably 2250 MPa or more.

On the other hand, if the drawn heat treated steel wire for highstrength spring use is too high in yield strength, the cold coilingability falls, so the yield strength is made 2400 MPa or less.

Yield strength (if yield strength cannot be seen, 0.2% proof stress):1600 to 1980 MPa

In the present invention, the yield strength or yield point of the drawnheat treated steel wire for high strength spring use means the top yieldstrength when a yield point is seen at the stress-strain curve in asingle-axis tensile test and the 0.2% proof stress when no yield pointis seen. To secure the strength or sag resistance of the spring, whichelastically deformed by repeated stress, raising the yield strength ispreferable. To raise the yield strength of the spring, raising the yieldstrength of the material, that is, the drawn heat treated steel wire forhigh strength spring use, is preferable.

On the other hand, if the drawn heat treated steel wire for highstrength spring use becomes high in yield strength, the cold coilingability is sometimes impaired. Therefore, the drawn heat treated steelwire for high strength spring use preferably has a yield strength of1600 MPa or more for securing the strength and sag property of thespring.

To impart further higher durability, 1700 MPa or more is preferable.

On the other hand, if the yield strength exceeds 1980 MPa, the coldcoiling ability is sometimes impaired, so the yield strength ispreferably made 1980 MPa or less. Note that to raise the yield strengthof the drawn heat treated steel wire for high strength spring use of thematerial having the same yield strength right after short time quenchingand tempering, it is preferable to lower the volume the volume rate ofthe residual austenite.

Vicker's hardness after nitriding by holding at 500° C. for 1 hour:Surface layer hardness HV≧750, internal hardness HV≧570

A high strength spring is improved in surface layer hardness innitriding, but the inside softens. For example, in gas soft nitriding at500° C., if the conventional heating temperature becomes 500° C., it wasdifficult to suppress softening at the inside of the drawn heat treatedsteel wire for high strength spring use. The drawn heat treated steelwire for high strength spring use of the present invention is excellentin temper softening resistance and enables fatigue characteristics andthe sag property of the spring after heating at 500° C. to be secured.In the present invention, the surface layer hardness and the internalhardness after gas soft nitriding are defined.

The surface layer hardness is made a micro Vicker's hardness at thedepth of 50 to 100 μm from the surface layer of 750 or more. If lessthan 750, the surface layer hardness becomes insufficient and thefatigue durability also becomes inferior, so residual stress after shotpeening cannot be sufficiently imparted. Preferably, the surface layerhardness is 780 or more.

On the other hand, in internal hardness, the Vicker's hardness issometimes measured when, in quenching, the temperature of the surfacelayer of the steel wire is higher than the inside, so measuring this ata position of 500 μm depth from the surface is preferable. To secure thespring fatigue characteristics and sag property, the Vicker's hardnessafter heat treatment holding the wire at 500° C. for 1 hour should be570 or more. Furthermore, 575 or more is preferable.

Note that, the upper limit of the Vicker's hardness after holding at500° C. for 1 hour for heat treatment is not particularly defined, butto ensure that the Vicker's hardness before the heat treatment is notexceeded, usually it is made 783 or less.

Furthermore, when using the drawn heat treated steel wire for highstrength spring use of the present invention as the material forproduction of high strength springs, the surface layer is hardened byshot peening, nitriding, etc. On the other hand, the Vicker's hardnessat a position of 500 μm depth from the surface of the high strengthspring (internal hardness) is affected by the heating in nitriding.Therefore, when actually producing a spring, the internal hardness willfluctuate depending on the temperature of the nitriding.

Note that, when using the drawn heat treated steel wire for highstrength spring use of the present invention as the material forproduction of high strength springs, it is cold coiled and nitrided. Forthis reason, the residual austenite at a position of 500 μm depth fromthe surface of the high strength springs falls somewhat compared withthe material of the drawn heat treated steel wire for high strengthspring use.

However, the chemical compositions, spherical carbides, and prioraustenite crystal grain size are believed to be little affected by thecold coiling and nitriding. Therefore, the chemical compositions,spherical carbides, and prior austenite crystal grain size of the highstrength steel made using the drawn heat treated steel wire for highstrength spring use of the present invention as a material are the sameextent as the chemical compositions, spherical carbides, and prioraustenite crystal grain size of the drawn heat treated steel wire forhigh strength spring use of the present invention.

Next, the method of production of the drawn heat treated steel wire forhigh strength spring use of the present invention will be explained.

A steel bloom adjusted to predetermined chemical compositions was rolledto produce a steel billet reduced in size. Further, the billet washeated, then hot rolled to obtain pre-drawn steel wire for high strengthspring use. This pre-drawn steel wire for high strength spring use waspatented, the shaped and, furthermore, was annealed for softening thehard layer. It was then drawn, quenched, and tempered to produce drawnheat treated steel wire for high strength spring use. The “patenting” isheat treatment for making the structure of the steel wire after hotrolling ferrite and pearlite and is performed for softening the steelwire before drawing. After drawing, oil tempering, induction hardeningtreatment, and other quenching and tempering are performed to adjust thesteel wire in structure and characteristics.

In the method of production of pre-drawn steel wire for high strengthspring use of the present invention, the process of preventingcoarsening of the spherical carbides is important.

In particular, when containing high C and high Cr like in the presentinvention, it is extremely important to sufficiently heat the bloom orbillet before rolling in that state and ease precipitation inside thesteel and to dissolve the internal coarse carbides (alloy carbides andcementite) and make the material uniform. To prevent the formation ofcoarse spherical carbides, the coarse carbides which are formed at thebloom or billet must be made to dissolve in the steel. Furthermore,causing uniform dispersion in the steel is necessary. For this reason,raising the heating temperature is preferable.

Therefore, first, the bloom or billet after casting is made a heatingtemperature of 1250° C. or more. Due to this, it is possible to make theundissolved spherical carbides sufficiently dissolve. For this reason,in the heating of the subsequent rolling, patenting, and quenching, theheating temperature and the heating time are insufficient, soundissolved spherical carbides easily remain, but to enable sufficientdissolution from the start, the dimensions of the undissolved sphericalcarbides can be controlled to less than 0.2 μm. The bloom heatingtemperature should be 1270° C. or more.

Next, the billet which is produced by rolling the bloom is further hotrolled (wire rod is rolled) to produce pre-drawn steel wire for highstrength spring use. At this time, the heating temperature of the billetis made 1200° C. or more. Preferably, the heating temperature of thebillet should be made 1250° C. or more.

After extracting the steel from the heating furnace, the temperaturefalls and precipitates grow. For this reason, after extraction from theheating furnace, the hot rolling is preferably completed within 5minutes. By the above heating of the bloom and billet, the coarsecarbides in the steel are uniformly dispersed and dissolved and canuniformly finely precipitate in the later precipitation.

Note that, when rolling a bloom into steel wire without going through abillet, the heating temperature before rolling of the bloom should bemade 1250° C. or more, more preferably 1270° C. or more.

In the above way, to suppress coarsening of the undissolved carbides ofthe steel wire after heat treatment, even if greatly reducing theundissolved carbides which are present before drawing (that is, afterwire rod rolling) and if for example undissolved carbides had beenpresent, it is necessary to make the size finer to prevent easycoarsening.

Therefore, in the rolling step of heating before drawing, it isimportant to make the bloom heating temperature and the billets heatingtemperature sufficiently high for the carbides to dissolve. Due to this,the size of the undissolved spherical carbides can be kept small. Therolling of the spring steel is completed in several minutes fromextraction of the billet from the heating furnace to a size of materialbefore drawing of about φ10 mm. For this reason, it is important to heatto 1200° C. or more where the effect of the billet heating temperatureis the largest. 1250° C. or more is more preferable. 1270° C. or more ismore preferable.

After rolling, the wire is taken up in a coil and air cooled at thattime as general practice. For this reason, usually the microstructure ofthe pre-drawn steel wire (steel wire after rolling of wire rod) iscomprised of ferrite and pearlite or pearlite with a high pearlitestructure fraction since the amount of C is high. Undissolved sphericalcarbides are present in the base material.

The undissolved spherical carbides can be observed by observing apolished and etched detection sample by an SEM. The undissolved carbidescan be clearly differentiated from the lamellar cementite contained inthe pearlite structure of the base material since they are spherical. Ofcourse, the magnitude may also be measured.

Due to the above step, a pre-drawn steel wire for spring use (rolledwire rod) is obtained.

After hot rolling, the pre-drawn steel wire for spring use is patented.The heating temperature of this patenting may be made 900° C. or more topromote dissolution of the carbides. A high temperature of 930° C. ormore is more preferable. Further, 950° C. or more is preferable. Afterthat, the wire may be patented at 600° C. or less. In the pre-drawnsteel wire for spring use according to the present invention, the methodof patenting and drawing is not limited. If a general patenting anddrawing method for steel wire, the same treatment as usual may beperformed.

When drawing by the wire diameter and precision required is omitted, thepatenting step before the drawing may be omitted. In this case, bymaking the heating temperature in the later explained quenching high(for example, 970° C. or more), dissolution of the undissolved sphericalcarbides is promoted.

The quenching after the drawing is performed by heating to temperatureof the A₃ point or more. To promote the dissolution of carbides, it ispreferable to raise the heating temperature of the quenching. In thequenching, to suppress the growth of carbides, the heating rate ispreferably made 10° C./sec or more and the holding time at thetemperature of the A₃ point or more is preferably made 1 minute to 5minutes. To suppress grain growth of the austenite, it is preferable toshorten the holding time. To promote the quenching and martensitetransformation, the cooling rate is preferably made 50° C./sec to 100°C.

The coolant in the quenching process is preferably made 100° C. or less,more preferably a low temperature of 80° C. or less, but in the presentinvention, to precisely control the amount of residual austenite, thecoolant temperature is made 40° C. or more. The coolant is notparticularly limited so long as being an oil, a water soluble quenchingagent, water, or other coolant which enables quenching. Further, thecooling time may be shortened like with oil tempering and inductionhardening treatment. It is preferable to avoid extending the holdingtime at a low temperature for greatly reducing the residual austeniteand lowering the coolant temperature to 30° C. or less. That is, thequenching is preferably ended within 5 minutes.

After quenching, tempering is performed. The tempering suppresses thegrowth of carbides, so it is preferable to make the heating rate 10°C./sec or more and make the holding time 15 minutes or less. The holdingtime fluctuates due to the chemical compositions and the targetedstrength, but the material is usually held at 400 to 500° C.

The pre-drawn steel wire for high strength spring use is cold coiled towork it to the desired spring shape, is relieved of stress, and isnitrided and shot peened to produce the spring.

The cold coiled steel wire is reheated by stress-relieving annealing,nitriding, etc. At this time, the inside is softened, so the performanceof the spring falls. In particular, in the present invention, even ifperforming the nitriding at a high temperature of about 500° C.,sufficient hardness is maintained. As a result, if using the pre-drawnsteel wire for high strength spring use of the present invention as amaterial, it is possible to make the micro Vicker's hardness at a depthof 500 μm from the surface layer of high strength springs HV575 or more.Note that, the micro Vicker's hardness is measured at a depth of 500 μmfrom the surface layer of the spring so as to evaluate the Vicker'shardness of the base material not affected by nitriding and shot peeningfor hardening.

EXAMPLES

Steels having the chemical compositions shown in Tables 1-1 to 1-4 weresmelted in a 10 kg vacuum melting furnace and cast to obtained blooms orbillets. These vacuum melted materials were hot forged up to φ8 mm.After that, the materials hot forged up to φ8 mm were heated at 1270°C.×4 hr. Further, part of the samples were refined in a 250 tonconverter, continuously cast to prepare blooms, then heated at 1270°C.×4 hr or more, then made into cross-section 160 mm×160 mm billets.Furthermore, these were rolled to obtained φ8 mm rolled wire rods. Theheating temperature of the billets before rolling was made 1200° C. ormore.

A diameter 8 mm pre-drawn steel wire (rolled wire rod) is preferablymade an easily drawn structure by patenting it before drawing. Theheating temperature at the patenting is preferably 900° C. or more sothat the carbides etc. sufficiently dissolve. The patenting is performedby heating at 930° C., then charging the sample into a 600° C. flowingbed. After patenting, the wire is drawn to obtain a diameter 4 mm drawnwire rod. In this way, by heating the bloom at a high temperature, thenmaking the temperature in the rolling process, patenting, and quenchingas high as possible, it is possible to suppress growth of undissolvedspherical carbides and keep the dimensions down to 0.2 μm or less.

To adjust the yield strength of the patented and drawn steel wire, thewire was quenched and tempered to produce pre-drawn steel wire forspring use. Note that, a sample which broke in the drawing was notquenched and tempered. The quenching and tempering were performed byheating the drawn steel wire by a 10° C./sec or more heating rate at950° C. or 1100° C. (temperature of A₃ point or more), holding at thepeak heating temperature for 4 minutes to 5 minutes, then placing thesteel in a room temperature water tank so that the cooling rate became50° C./sec or more and cooling down to 100° C. or less.

As the results of evaluation, the state of wire breakage, prioraustenite grain size number, residual austenite amount (vol %), circleequivalent diameter and density of presence of carbides, yield strength,0.2% proof stress, notch bending angle, average fatigue strength, andVicker's hardness after gas soft nitriding are shown.

The target values to be passed were made as follows with reference toconventional steel wire for high strength spring use.

Prior austenite grain size number: 10 degrees or more

Residual austenite amount (vol %): 20% or less

Circle equivalent diameter of spherical carbides: 0.2 μm or less

Yield strength: 2100 MPa or more

0.2% proof stress: 1800 MPa or more

Yield ratio: 75% to 95%

Notch bending angle: 28 degrees or more

Average fatigue strength (Nakamura type rotating bending strength): 900MPa or more

Internal hardness by Vicker's hardness after gas nitriding: 590 Hv ormore

Nitrided layer hardness by Vicker's hardness after gas nitriding: 750 Hvor more

Note that, in the steel wire according to the present invention, thestrength and workability (coiling ability) both have to be achieved, soif the yield ratio is too high, the workability deteriorates. Therefore,the upper limit of the yield ratio is preferably 90%, more preferably88% or less.

A sample was taken from the obtained drawn heat treated steel wire forspring use, evaluated for prior austenite grain size, volume rate ofresidual austenite, and carbides, then was subjected to a tensile test,notch bending test, and micro Vicker's hardness test. Note that, thefatigue characteristics were evaluated by treatment simulatingproduction of a spring (below, referred to as “spring production andtreatment”) including gas nitriding simulating nitriding performed onthe spring after working (500° C., 60 minutes), shot peening (diameterof cut wire 0.6 mm, 20 minutes), and low temperature stress-relievingtreatment (180° C., 20 minutes).

The prior austenite grain size number was measured based on JIS G 0551.The circle equivalent diameter and density of presence of the carbideswere measured by using an electrolytically etched sample, obtaining aSEM structural photograph, and analyzing the image. Further, the volumerate of the residual austenite was measured by the magnetic measurementmethod.

The fatigue test is a Nakamura type rotating bending fatigue test(fatigue test bending by two-point supported weight and turning by motorto apply compressive and tensile stress to surface of wire). The maximumload force of 10 samples showing a lifetime of 10⁷ cycles or more by aprobability of 50% or more was made the average fatigue strength. Thenotch bending test is a test for evaluating the cold coiling ability andis performed as follows.

A punch 2 with an angle of the tip shown in FIG. 2 of 120° was used toprovide a groove (notch) of a maximum depth of 30 μm in the test piece.Note that, as shown in FIG. 3, the notch 4 was provided at a right angleto the longitudinal direction at the center of the test piece 3 in thelongitudinal direction. Next, as shown in FIG. 4, from the opposite sideof the notch 4, a pusher 5 was used to apply a load P of a maximumtensile stress through a load-use fixture 6 and the test piece wasdeformed by three-point bending. Note that, the radius of curvature r ofthe tip of the load-use fixture 6 was made 4.0 mm, while the differenceL between supports was made L=2r+3D. Here, D is the diameter of the testpiece.

The bending deformation continued to be applied until the notch partfractured. The bending angle at the time of fracture (notch bendingangle) was measured as shown in FIG. 5. Note that, when the test piecewas split, the fractured parts were placed together to measure the notchbending angle θ. In the present invention, a sample with a notch bendingangle of 28° or more is judged to be excellent in cold coiling ability.

The micro Vicker's hardness after nitriding was evaluated using thedepth of 500 μm or more from the surface layer as the internal hardnesswas defining the micro Vicker's hardness of a depth of 50 μm from thesurface layer as the “nitrided layer hardness”. The measurement weightwas 10 g.

The results of these tests are shown in Tables 1-5 to 1-8. Note that, inTables 1-5 to 1-8, the metal structure is comprised of residualaustenite (γ) plus tempered martensite and slight inclusions. Further,the balance of the chemical compositions was iron and unavoidableimpurities.

The pre-drawn steel wire (steel wire after rolling wire rod) wasevaluated only by the circle equivalent diameter of the undissolvedspherical carbides. This is because since this is before heat treatment,even if measuring the mechanical properties or the austenite grain sizeetc., there is not much meaning to the figures.

Examples 1 to 47 of the present invention all have the indicator of thecold coiling ability, that is, the notch bending angle, of a good 28° ormore and have an excellent indicator of the spring durability, that is,the Nakamura type rotating bending fatigue strength (hereinafter simplyreferred to as the “fatigue durability”) and an excellent indicator ofthe sag property and temper softening resistance, that is, the nitridedlayer hardness.

Comparative Examples 48 and 49 are examples where the amount of additionof C is outside the range of the claims. If C is over the prescribedamount (Comparative Example 48), the undissolved spherical carbidesbecome greater and the indicator of the cold coiling ability, the notchbending angle, is low. On the other hand, if C is smaller than theprescribed amount (Comparative Example 49), a sufficient yield strengthcannot be secured. In particular, the internal hardness after nitridingbecomes lower and the spring fatigue durability (Nakamura type rotatingbending fatigue strength) and the sag property (internal hardness afternitriding).

Comparative Examples 50 and 51 are examples where the amount of additionof Si is outside the range of the claims. If Si exceeds the prescribedamount, the matrix is embrittled and the workability is impaired, thatis, the notch bending angle is low. On the other hand, if Si is smallerthan the prescribed amount, the quenching and tempering characteristicsdeteriorate, so sufficient strength cannot be secured after heating bynitriding. In particular, the internal hardness after nitriding and thenitrided layer hardness become low.

Comparative Examples 52 and 53 are examples where the amount of additionof Mn is outside the range of the claims. If Mn is over the prescribedrange, the residual austenite becomes greater, the yield strength falls,and the fatigue durability (Nakamura type rotating bending fatiguestrength) is inferior. On the other hand, when Mn is smaller than theprescribed amount, the residual austenite falls too much and theworkability deteriorates, so the notch bending angle falls.

Comparative Examples 54 and 55 are examples where the amount of additionof Cr is outside the range of the claims. If the Cr is over theprescribed range, cementite stabilizes and even in the high temperatureheating of the bloom or billet, quenching and tempering, etc.,undissolved carbides increase and the spring workability is greatlyreduced. For this reason, the notch bending angle falls. On the otherhand, if Cr is smaller than the prescribed amount, the steel ends upsoftening in the heat treatment in the nitriding etc. and the so-calledtemper softening resistance otherwise becomes insufficiently so thenitrided layer hardness falls.

Comparative Examples 56, 57, and 58 are examples where the amounts ofaddition of Mo, W, and Mo+W are over the ranges of the claims. If Mo andW exceed the prescribed amounts, in rolling and cooling and afterpatenting and other heat treatment, a supercooled structure ofmartensite, bainite, etc. forms, the wire breaks in the conveyance ordrawing process, and the measurement test cannot be performed.

Comparative Example 59 is an example of excessive addition of V. V is anelement which forms carbides in the steel. Excessive addition causesundissolved carbides to form around the V, the workability todeteriorate, and the notch bending angle to fall.

Comparative Examples 60 and 61 are cases where the amount of N isexcessive compared with the range of the claims. This excessive N raisesthe temperature of formation of nitrides and carbonitrides of V, Nb,etc. and causes coarsening of carbides and other precipitates usingthese as nuclei. Further, when used for repeated heating such as in thepresent invention, the nitrides, carbonitride, and carbides areincompletely dissolved and a large amount of coarse undissolvedspherical carbides remain. As a result, the workability is impaired.This is an example where the notch bending angle falls.

Comparative Examples 62 and 63 are examples where the amount of additionof Nb is outside the range of the claims. If Nb exceeds the prescribedamount, the hot ductility is remarkably impaired, numerous surface flawsoccur at the rolled material, wire breakage occurs during drawing, and ameasurement test could not be run.

Comparative Examples 64 is the case where the sum of the amounts ofaddition of Mn and V is more than the range explained in the presentinvention. The amount of residual austenite in the steel wire becomesgreater than the prescribed value. In the notch bending test, the notchpart hardens due to the stress-induced transformation and theworkability falls. This is an example where the notch bending anglefalls. While repeating ourselves, V is not added in the presentinvention, but sometimes V is included as an unavoidable impurity, sothis is a limitation for rendering the V harmless.

Comparative Examples 65 is the case where the sum of the amounts ofaddition of Mn and V is lower than the range explained in the presentinvention. The amount of residual austenite is smaller than the optimumrange, so the workability, that is, the notch bending angle, falls.

Comparative Example 66 is the case where the sum of the amounts ofaddition of Cr and V is greater than the scope explained in the presentinvention. The undissolved spherical carbides excessively remain and theworkability, that is, the notch bending angle, falls.

Comparative Example 67 is the case where the sum of the amounts ofaddition of Cr and V is less than the range explained in the presentinvention. The workability is excellent, but the internal hardness afternitriding and the nitrided layer hardness are insufficient and thespring performance is not sufficient.

Comparative Examples 68 to 70 are cases where the difference between theamount of Si and the amount of Cr ([Si %]-[Cr %]) is off from the scopeof the claims and the amount of Cr is greater than the amount of Si. IfCr is excessive with respect to the amount of Si, undissolved sphericalcarbides remain and the workability is degraded, that is, that is, thenotch bending angle falls.

Similarly, Comparative Examples 71 and 72 are the case where thedifference of the amount of Si and the amount of Cr ([Si %]-[Cr %]) islarger than the upper limit of the range of the claims. Si is veryexcessive compared with the amount of Cr. In these cases, the surfacelayer decarburized layer of the rolled material greatly grows and cannotbe sufficiently removed by a slight amount of surface layer shaving. Forthis reason, the fatigue durability (Nakamura type rotating bendingfatigue strength) was inferior.

Comparative Examples 73 and 74 are respectively the Invention Example 1and Invention Example 23 where the steel is rolled at the billet heatingtemperature 1100° C. At the start of the rolling, undissolved sphericalcarbides remain. The effects finally remain, so the workability isdegraded, that is, the notch bending angle falls.

Invention Examples 101 to 109 are examples of the pre-drawn steel wiresof Invention Examples 1 to 5 and 20 to 23. Comparative Examples 110 and111 are the Invention Examples 101 and 106 where the billet heatingtemperature is made 1100° C.

The pre-drawn steel wire is evaluated, so only the maximum circleequivalent diameter of the undissolved spherical carbides is evaluated.If the billet heating temperature is high, it is learned that the circleequivalent diameter of the undissolved spherical carbides becomessmaller.

TABLE 1-1 Chemical compositions (mass %) Ex. C Si Mn P S Cr Al N V Nb  1Inv. ex. 0.78 2.48 0.68 0.0076 0.0045 1.57 0.0022 0.0031 — —  2 Inv. ex.0.77 2.41 0.68 0.0034 0.0047 2.05 0.0011 0.0042 — —  3 Inv. ex. 0.682.38 0.87 0.0045 0.0061 1.53 0.0013 0.0033 — —  4 Inv. ex. 0.88 2.500.87 0.0063 0.0071 1.71 0.0018 0.0035 — —  5 Inv. ex. 0.78 2.11 0.840.0057 0.0039 1.50 0.0010 0.0031 — —  6 Inv. ex. 0.72 2.62 0.62 0.00540.0031 1.96 0.0022 0.0038 — —  7 Inv. ex. 0.72 2.67 0.58 0.0037 0.00351.51 0.0015 0.0032 — —  8 Inv. ex. 0.76 2.28 1.02 0.0067 0.0069 1.820.0028 0.0031 — —  9 Inv. ex. 0.73 2.23 0.73 0.0054 0.0077 1.53 0.00150.0058 — — 10 Inv. ex. 0.77 2.35 0.80 0.0041 0.0047 1.52 0.0019 0.0032 —— 11 Inv. ex. 0.77 2.59 0.83 0.0060 0.0074 1.53 0.0012 0.0063 — 0.008 12Inv. ex. 0.75 2.52 0.68 0.0054 0.0056 1.67 0.0013 0.0039 0.06 — 13 Inv.ex. 0.75 2.33 0.83 0.0066 0.0060 2.00 0.0030 0.0037 0.09 0.001 14 Inv.ex. 0.72 2.41 0.86 0.0040 0.0068 1.77 0.0013 0.0033 — — 15 Inv. ex. 0.772.57 0.75 0.0078 0.0075 1.91 0.0018 0.0043 — — 16 Inv. ex. 0.72 2.420.84 0.0044 0.0053 1.90 0.0019 0.0056 — — 17 Inv. ex. 0.76 2.53 0.670.0061 0.0076 1.65 0.0028 0.0039 — — 18 Inv. ex. 0.73 2.46 0.66 0.00510.0060 1.60 0.0023 0.0035 — — 19 Inv. ex. 0.73 2.34 0.75 0.0039 0.00711.97 0.0028 0.0032 — — 20 Inv. ex. 0.73 2.35 0.70 0.0046 0.0033 1.910.0016 0.0034 0.10 0.008 21 Inv. ex. 0.77 2.46 0.71 0.0031 0.0054 1.950.0013 0.0038 0.03 0.003 22 Inv. ex. 0.76 2.35 0.64 0.0051 0.0072 1.720.0016 0.0044 0.07 0.007 23 Inv. ex. 0.73 2.36 0.73 0.0033 0.0073 1.800.0019 0.0033 0.08 0.006 24 Inv. ex. 0.76 3.20 0.72 0.0076 0.0038 2.100.0010 0.0037 0.06 — Chemical compositions (mass %) Ex. Mo W Mg Zr CaMn + V Cr + V Si—Cr Mo + W  1 Inv. ex. — — — — — — — 0.90 —  2 Inv. ex.— — — — — — — 0.35 —  3 Inv. ex. — — — — — — — 0.85 —  4 Inv. ex. — — —— — — — 0.79 —  5 Inv. ex. — — — — — — — 0.61 —  6 Inv. ex. — — — — — —— 0.66 —  7 Inv. ex. — — — — — — — 1.16 —  8 Inv. ex. — — — — — — — 0.46—  9 Inv. ex. — — — — — — — 0.70 — 10 Inv. ex. — — — — — — — 0.83 — 11Inv. ex. — — — — — — — 1.07 — 12 Inv. ex. — — — — — 0.74 1.74 0.84 — 13Inv. ex. — — — — — 0.93 2.10 0.32 — 14 Inv. ex. 0.12 — — — — — — 0.640.12 15 Inv. ex. — 0.15 — — — — — 0.66 0.15 16 Inv. ex. 0.12 0.16 — — —— — 0.52 0.27 17 Inv. ex. — — 0.0005 — — — — 0.87 — 18 Inv. ex. — — —0.0002 — — — 0.86 — 19 Inv. ex. — — — — 0.0011 — — 0.37 — 20 Inv. ex.0.12 0.17 0.0002 0.0003 0.0011 0.79 2.00 0.44 0.28 21 Inv. ex. 0.15 0.160.0003 0.0002 0.0003 0.74 1.98 0.51 0.31 22 Inv. ex. 0.11 0.17 0.00050.0003 0.0006 0.71 1.79 0.64 0.28 23 Inv. ex. 0.11 0.15 0.0004 0.00010.0012 0.81 1.88 0.56 0.27 24 Inv. ex. 0.10 — — — — 0.77 2.16 1.10 0.10

TABLE 1-2 Chemical compositions (mass %) Ex. C Si Mn P S Cr Al N V Nb 25Inv. ex. 0.73 2.05 0.80 0.0062 0.0052 1.71 0.0011 0.0033 0.08 0.009 26Inv. ex. 0.76 2.52 0.71 0.0069 0.0080 1.69 0.0024 0.0039 0.05 — 27 Inv.ex. 0.75 2.30 0.82 0.0074 0.0055 1.73 0.0018 0.0038 0.08 0.010 28 Inv.ex. 0.73 2.24 1.20 0.0052 0.0074 1.52 0.0027 0.0033 0.09 — 29 Inv. ex.0.73 2.44 0.74 0.0076 0.0046 1.73 0.0024 0.0032 0.05 0.010 30 Inv. ex.0.74 2.24 0.76 0.0052 0.0076 1.41 0.0016 0.0030 0.05 0.005 31 Inv. ex.0.75 2.28 0.89 0.0065 0.0043 1.70 0.0011 0.0043 0.03 0.010 32 Inv. ex.0.74 2.23 0.89 0.0043 0.0049 1.50 0.0022 0.0032 0.04 — 33 Inv. ex. 0.772.86 0.77 0.0057 0.0057 2.48 0.0015 0.0040 0.09 0.005 34 Inv. ex. 0.772.29 0.79 0.0055 0.0072 1.85 0.0016 0.0032 0.09 0.004 35 Inv. ex. 0.732.52 0.76 0.0078 0.0067 2.00 0.0022 0.0036 0.09 0.009 36 Inv. ex. 0.772.31 0.89 0.0046 0.0080 1.91 0.0021 0.0032 0.06 0.000 37 Inv. ex. 0.732.53 0.69 0.0046 0.0060 2.01 0.0022 0.0042 — 0.005 38 Inv. ex. 0.76 2.350.81 0.0034 0.0036 1.80 0.0016 0.0040 0.04 0.006 39 Inv. ex. 0.74 2.380.72 0.0031 0.0064 1.82 0.0010 0.0043 0.07 0.009 40 Inv. ex. 0.75 2.320.72 0.0034 0.0038 1.56 0.0017 0.0033 0.09 0.002 41 Inv. ex. 0.76 2.370.69 0.0056 0.0067 1.63 0.0028 0.0041 0.04 — 42 Inv. ex. 0.76 2.48 0.730.0050 0.0035 1.97 0.0016 0.0034 0.09 — 43 Inv. ex. 0.72 2.26 0.680.0053 0.0043 1.51 0.0027 0.0032 0.05 — 44 Inv. ex. 0.76 2.38 0.860.0077 0.0039 1.61 0.0020 0.0052 — — 45 Inv. ex. 0.77 2.23 0.76 0.00600.0061 1.69 0.0029 0.0053 0.05 0.006 46 Inv. ex. 0.76 2.35 0.86 0.00600.0067 1.87 0.0025 0.0035 0.05 0.010 47 Inv. ex. 0.71 2.28 0.77 0.00500.0058 1.88 0.0030 0.0033 0.07 0.003 Chemical compositions (mass %) Ex.Mo W Mg Zr Ca Mn + V Cr + V Si—Cr Mo + W 25 Inv. ex. 0.15 0.16 — — —0.88 1.79 0.34 0.31 26 Inv. ex. 0.28 — 0.0002 0.0003 0.0015 0.76 1.740.84 0.28 27 Inv. ex. 0.13 0.15 0.0004 0.0001 0.0005 0.91 1.82 0.56 0.2828 Inv. ex. 0.11 — 0.0003 0.0001 0.0011 1.29 1.62 0.72 0.11 29 Inv. ex.0.12 0.15 0.0001 0.0003 0.0013 0.79 1.78 0.71 0.27 30 Inv. ex. 0.11 0.160.0003 0.0001 0.0011 0.81 1.46 0.83 0.27 31 Inv. ex. 0.14 0.16 0.00040.0002 0.0012 0.93 1.73 0.58 0.31 32 Inv. ex. 0.11 — 0.0001 0.00010.0002 0.93 1.54 0.73 0.11 33 Inv. ex. 0.11 0.16 0.0002 0.0002 0.00050.86 2.57 0.38 0.27 34 Inv. ex. 0.23 0.15 0.0004 0.0003 0.0006 0.88 1.930.45 0.38 35 Inv. ex. 0.18 0.28 0.0005 0.0001 0.0005 0.84 2.09 0.52 0.4636 Inv. ex. 0.13 0.15 0.0001 0.0002 0.0006 0.95 1.97 0.40 0.28 37 Inv.ex. 0.14 0.17 0.0005 0.0002 0.0006 — — 0.52 0.31 38 Inv. ex. 0.12 0.090.0004 0.0001 0.0009 0.85 1.84 0.55 0.21 39 Inv. ex. 0.11 0.28 0.00020.0002 0.0007 0.80 1.89 0.55 0.39 40 Inv. ex. 0.13 0.15 0.0004 0.00020.0012 0.81 1.65 0.76 0.28 41 Inv. ex. 0.14 0.17 0.0002 0.0002 0.00080.73 1.67 0.74 0.30 42 Inv. ex. 0.10 0.16 — — — 0.82 2.06 0.51 0.26 43Inv. ex. 0.14 0.14 0.0002 0.0003 0.0002 0.73 1.56 0.75 0.29 44 Inv. ex.0.14 0.15 — — — — — 0.77 0.28 45 Inv. ex. 0.11 0.15 — — — 0.81 1.74 0.540.26 46 Inv. ex. 0.10 0.15 0.0005 0.0003 0.0010 0.90 1.92 0.48 0.25 47Inv. ex. 0.10 0.17 0.0005 0.0003 0.0006 0.84 1.95 0.41 0.27

TABLE 1-3 Chemical compositions (mass %) Ex. C Si Mn P S Cr Al N V Nb 48Comp. ex. 0.95 2.47 0.61 0.0033 0.0078 1.54 0.0022 0.0043 0.09 0.006 49Comp. ex. 0.58 2.59 0.62 0.0079 0.0057 1.81 0.0021 0.0066 0.09 0.005 50Comp. ex. 0.71 3.80 0.85 0.0049 0.0035 2.03 0.0017 0.0032 0.05 0.002 51Comp. ex. 0.73 1.86 0.83 0.0034 0.0060 1.32 0.0013 0.0031 0.09 0.004 52Comp. ex. 0.72 2.46 1.54 0.0076 0.0066 2.02 0.0024 0.0043 0.06 0.008 53Comp. ex. 0.72 2.22 0.21 0.0057 0.0064 1.58 0.0024 0.0036 0.09 0.005 54Comp. ex. 0.73 3.13 0.71 0.0048 0.0030 2.72 0.0014 0.0030 0.04 0.007 55Comp. ex. 0.78 2.21 0.68 0.0051 0.0041 1.02 0.0013 0.0031 0.09 0.011 56Comp. ex. 0.75 2.42 0.74 0.0034 0.0067 1.78 0.0017 0.0033 0.09 0.005 57Comp. ex. 0.73 2.45 0.81 0.0046 0.0061 1.75 0.0027 0.0046 0.11 0.010 58Comp. ex. 0.73 2.30 0.64 0.0044 0.0034 1.80 0.0011 0.0037 — 0.001 59Comp. ex. 0.74 2.23 0.80 0.0060 0.0046 1.72 0.0026 0.0032 0.46 0.007 60Comp. ex. 0.77 2.43 0.83 0.0062 0.0072 1.96 0.0025 0.0076 0.08 — 61Comp. ex. 0.75 2.26 0.78 0.0069 0.0058 1.99 0.0013 0.0085 0.07 0.010 62Comp. ex. 0.73 2.50 0.62 0.0049 0.0062 1.73 0.0020 0.0053 0.05 0.035 63Comp. ex. 0.72 2.21 0.72 0.0074 0.0051 1.86 0.0026 0.0036 — 0.024 64Comp. ex. 0.74 2.50 1.18 0.0053 0.0056 2.02 0.0022 0.0042 0.09 0.006 65Comp. ex. 0.74 2.52 0.51 0.0064 0.0046 1.73 0.0019 0.0039 0.06 0.003 66Comp. ex. 0.75 2.76 0.72 0.0048 0.0058 2.45 0.0014 0.0032 0.09 0.004 67Comp. ex. 0.77 2.20 0.79 0.0059 0.0059 1.31 0.0016 0.0035 0.03 0.002 68Comp. ex. 0.76 2.12 0.66 0.0068 0.0075 2.31 0.0026 0.0033 0.09 — 69Comp. ex. 0.74 2.10 0.88 0.0056 0.0056 2.23 0.0023 0.0043 0.06 0.005 70Comp. ex. 0.78 2.23 0.73 0.0067 0.0075 2.41 0.0023 0.0048 0.11 0.000 71Comp. ex. 0.71 3.12 0.76 0.0068 0.0044 1.54 0.0021 0.0035 0.07 0.005 72Comp. ex. 0.76 3.45 0.66 0.0071 0.0052 1.66 0.0017 0.0036 0.04 0.005 73Comp. ex. 0.78 2.48 0.68 0.0076 0.0045 1.57 0.0022 0.0031 — — 74 Comp.ex. 0.73 2.36 0.73 0.0033 0.0073 1.80 0.0019 0.0033 0.08 0.006 Chemicalcompositions (mass %) Ex. Mo W Mg Zr Ca Mn + V Cr + V Si—Cr Mo + W 48Comp. ex. 0.12 0.14 0.0004 0.0001 0.0013 0.70 1.63 0.93 0.26 49 Comp.ex. 0.15 0.15 0.0003 0.0001 0.0014 0.71 1.90 0.77 0.29 50 Comp. ex. 0.140.16 0.0002 0.0003 0.0008 0.90 2.09 1.77 0.31 51 Comp. ex. 0.14 0.140.0004 0.0003 0.0005 0.92 1.41 0.54 0.29 52 Comp. ex. 0.14 0.15 0.00030.0003 0.0014 1.60 2.08 0.44 0.29 53 Comp. ex. 0.15 0.17 0.0001 0.00020.0011 0.30 1.67 0.64 0.31 54 Comp. ex. 0.14 0.15 0.0003 0.0002 0.00110.75 2.76 0.41 0.29 55 Comp. ex. 0.14 0.17 0.0002 0.0001 0.0012 0.771.11 1.19 0.31 56 Comp. ex. 0.42 0.07 0.0004 0.0003 0.0005 0.83 1.870.63 0.49 57 Comp. ex. 0.12 0.50 0.0003 0.0001 0.0008 0.92 1.86 0.700.62 58 Comp. ex. 0.26 0.27 0.0005 0.0002 0.0007 — — 0.50 0.53 59 Comp.ex. 0.11 0.16 0.0004 0.0001 0.0015 1.26 2.18 0.51 0.27 60 Comp. ex. 0.150.17 0.0002 0.0001 0.0004 0.91 2.04 0.47 0.32 61 Comp. ex. 0.10 0.160.0003 0.0002 0.0005 0.85 2.06 0.27 0.27 62 Comp. ex. 0.13 0.17 — — —0.67 1.78 0.77 0.29 63 Comp. ex. 0.12 0.17 0.0002 0.0001 0.0012 — — 0.340.29 64 Comp. ex. 0.10 0.16 — — — 1.27 2.11 0.48 0.26 65 Comp. ex. 0.120.16 0.0004 0.0002 0.0007 0.57 1.79 0.79 0.28 66 Comp. ex. 0.11 0.150.0002 0.0002 0.0015 0.81 2.54 0.31 0.26 67 Comp. ex. 0.14 0.16 0.00030.0001 0.0011 0.82 1.34 0.89 0.30 68 Comp. ex. 0.15 0.16 0.0003 0.00010.0013 0.75 2.40 −0.19 0.31 69 Comp. ex. 0.14 0.16 — — — 0.94 2.29 −0.130.30 70 Comp. ex. 0.12 0.17 0.0004 0.0003 0.0013 0.84 2.52 −0.18 0.28 71Comp. ex. 0.12 0.15 — — — 0.84 1.61 1.58 0.26 72 Comp. ex. 0.11 0.160.0004 0.0002 0.0002 0.70 1.70 1.79 0.26 73 Comp. ex. — — — — — — — 0.90— 74 Comp. ex. 0.11 0.15 0.0004 0.0001 0.0012 0.81 1.88 0.56 0.27

TABLE 1-4 Chemical compositions (mass %) Ex. C Si Mn P S Cr Al N V Nb101 Inv. ex. 0.78 2.48 0.68 0.0076 0.0045 1.57 0.0022 0.0031 — — 102Inv. ex. 0.77 2.41 0.68 0.0034 0.0047 2.05 0.0011 0.0042 — — 103 Inv.ex. 0.68 2.38 0.87 0.0045 0.0061 1.53 0.0013 0.0033 — — 104 Inv. ex.0.88 2.50 0.87 0.0063 0.0071 1.71 0.0018 0.0035 — — 105 Inv. ex. 0.782.11 0.84 0.0057 0.0039 1.50 0.0010 0.0031 — — 106 Inv. ex. 0.73 2.350.70 0.0046 0.0033 1.91 0.0016 0.0034 0.10 0.008 107 Inv. ex. 0.77 2.460.71 0.0031 0.0054 1.95 0.0013 0.0038 0.03 0.003 108 Inv. ex. 0.76 2.350.64 0.0051 0.0072 1.72 0.0016 0.0044 0.07 0.007 109 Inv. ex. 0.73 2.360.73 0.0033 0.0073 1.80 0.0019 0.0033 0.08 0.006 110 Comp. ex. 0.78 2.480.68 0.0076 0.0045 1.57 0.0022 0.0031 — — 111 Comp. ex. 0.73 2.36 0.730.0033 0.0073 1.80 0.0019 0.0033 0.08 0.006 Chemical compositions (mass%) Ex. Mo W Mg Zr Ca Mn + V Cr + V Si—Cr Mo + W 101 Inv. ex. — — — — — —— 0.90 — 102 Inv. ex. — — — — — — — 0.35 — 103 Inv. ex. — — — — — — —0.85 — 104 Inv. ex. — — — — — — — 0.79 — 105 Inv. ex. — — — — — — — 0.61— 106 Inv. ex. 0.12 0.17 0.0002 0.0003 0.0011 0.79 2.00 0.44 0.28 107Inv. ex. 0.15 0.16 0.0003 0.0002 0.0003 0.74 1.98 0.51 0.31 108 Inv. ex.0.11 0.17 0.0005 0.0003 0.0006 0.71 1.79 0.64 0.28 109 Inv. ex. 0.110.15 0.0004 0.0001 0.0012 0.81 1.88 0.56 0.27 110 Comp. ex. — — — — — —— 0.90 — 111 Comp. ex. 0.11 0.15 0.0004 0.0001 0.0012 0.81 1.88 0.560.27

TABLE 1-5 Billet Patenting Quenching Wire break etc. Max. Prior heatingtemp. heating temp. heating temp. [good = no spherical carbide austenitegrain Residual austenite Ex. (° C.) (° C.) (° C.) abnormality] diameterμm size (γ#) (vol %)  1 Inv. ex. 1250 930 950 good 0.06 10 9  2 Inv. ex.1250 930 950 good 0.13 12 9  3 Inv. ex. 1250 930 950 good 0.06 11 8  4Inv. ex. 1250 930 950 good 0.09 12 13  5 Inv. ex. 1250 930 950 good 0.1313 11  6 Inv. ex. 1250 930 950 good 0.10 12 10  7 Inv. ex. 1200 930 950good 0.08 10 10  8 Inv. ex. 1250 930 950 good 0.02 10 6  9 Inv. ex. 1250930 950 good 0.07 13 11 10 Inv. ex. 1200 930 1010 good 0.11 13 9 11 Inv.ex. 1250 930 1010 good 0.13 10 10 12 Inv. ex. 1250 930 1010 good 0.10 127 13 Inv. ex. 1250 930 950 good 0.10 13 13 14 Inv. ex. 1250 930 950 good0.07 13 10 15 Inv. ex. 1250 930 950 good 0.10 11 12 16 Inv. ex. 1250 930950 good 0.01 12 9 17 Inv. ex. 1250 930 950 good 0.12 11 13 18 Inv. ex.1250 930 950 good 0.09 12 9 19 Inv. ex. 1250 930 950 good 0.11 10 7 20Inv. ex. 1250 930 950 good 0.09 10 8 21 Inv. ex. 1250 930 950 good 0.0412 12 22 Inv. ex. 1250 — 970 good 0.04 10 12 23 Inv. ex. 1250 930 950good 0.06 12 11 24 Inv. ex. 1200 930 950 good 0.10 11 11 Notch InternalHardness of Tensile bending deg. Nakamura type rotary hardness afternitrided Ex. strength (MPa) 0.2% proof stress Yield ratio (%) (deg.)bending (MPa) nitriding (HV) layer (HV)  1 Inv. ex. 2214 1887 85 36 927597 788  2 Inv. ex. 2288 1939 85 39 918 638 787  3 Inv. ex. 2383 1945 8237 924 599 819  4 Inv. ex. 2217 1811 82 39 915 621 799  5 Inv. ex. 23051844 80 36 918 627 817  6 Inv. ex. 2241 1899 85 39 918 623 816  7 Inv.ex. 2243 1978 88 39 911 618 793  8 Inv. ex. 2301 1811 79 38 914 600 789 9 Inv. ex. 2293 1880 82 35 928 607 806 10 Inv. ex. 2263 1801 80 37 912612 795 11 Inv. ex. 2283 1939 85 37 929 608 808 12 Inv. ex. 2330 1844 7938 913 616 780 13 Inv. ex. 2169 1904 88 39 925 627 819 14 Inv. ex. 23461854 79 35 919 612 808 15 Inv. ex. 2189 1917 88 36 924 599 804 16 Inv.ex. 2285 1822 80 38 916 625 805 17 Inv. ex. 2248 1950 87 40 929 595 78718 Inv. ex. 2193 1904 87 38 925 616 804 19 Inv. ex. 2180 1949 89 36 928598 786 20 Inv. ex. 2374 1891 80 39 924 612 817 21 Inv. ex. 2368 1846 7836 911 638 806 22 Inv. ex. 2254 1853 82 40 914 629 795 23 Inv. ex. 23111852 80 39 923 625 783 24 Inv. ex. 2250 1956 87 35 911 616 793

TABLE 1-6 Billet Patenting Quenching Wire break etc. Max. Prior heatingtemp. heating temp. heating temp. [good = no spherical carbide austenitegrain Residual austenite Ex. (° C.) (° C.) (° C.) abnormality] diameterμm size (γ#) (vol %) 25 Inv. ex. 1250 930 950 good 0.10 10 8 26 Inv. ex.1250 930 950 good 0.12 11 13 27 Inv. ex. 1250 930 950 good 0.11 11 12 28Inv. ex. 1200 930 950 good 0.05 11 13 29 Inv. ex. 1250 930 950 good 0.0913 9 30 Inv. ex. 1250 930 950 good 0.04 11 7 31 Inv. ex. 1250 930 950good 0.07 12 10 32 Inv. ex. 1250 930 950 good 0.01 11 12 33 Inv. ex.1200 930 950 good 0.02 10 10 34 Inv. ex. 1250 930 950 good 0.03 11 11 35Inv. ex. 1250 930 950 good 0.01 12 9 36 Inv. ex. 1250 930 950 good 0.0512 13 37 Inv. ex. 1250 — 970 good 0.01 10 9 38 Inv. ex. 1250 930 950good 0.07 11 6 39 Inv. ex. 1250 930 950 good 0.15 12 6 40 Inv. ex. 1250930 950 good 0.06 13 8 41 Inv. ex. 1250 930 950 good 0.01 10 8 42 Inv.ex. 1250 930 950 good 0.06 13 6 43 Inv. ex. 1250 930 950 good 0.08 12 1244 Inv. ex. 1250 930 950 good 0.06 12 8 45 Inv. ex. 1250 930 950 good0.09 12 6 46 Inv. ex. 1250 930 950 good 0.13 12 12 47 Inv. ex. 1250 930950 good 0.03 11 7 Notch Internal Hardness of Tensile bending deg.Nakamura type rotary hardness after nitrided Ex. strength (MPa) 0.2%proof stress Yield ratio (%) (deg.) bending (MPa) nitriding (HV) layer(HV) 25 Inv. ex. 2357 1953 83 39 916 619 792 26 Inv. ex. 2195 1865 85 35930 628 794 27 Inv. ex. 2226 1872 84 38 922 592 795 28 Inv. ex. 21811854 85 38 910 631 784 29 Inv. ex. 2277 1890 83 37 916 635 786 30 Inv.ex. 2330 1938 83 37 922 593 798 31 Inv. ex. 2305 1867 81 38 921 630 79132 Inv. ex. 2216 1880 85 35 919 592 813 33 Inv. ex. 2285 1929 84 39 911597 797 34 Inv. ex. 2329 1944 83 39 916 601 794 35 Inv. ex. 2310 1824 7935 911 609 793 36 Inv. ex. 2153 1951 91 37 927 615 810 37 Inv. ex. 22201904 86 35 919 606 816 38 Inv. ex. 2178 1971 91 38 921 605 803 39 Inv.ex. 2268 1820 80 39 920 590 784 40 Inv. ex. 2302 1860 81 39 919 604 81441 Inv. ex. 2190 1896 87 37 924 611 806 42 Inv. ex. 2218 1893 85 37 927639 787 43 Inv. ex. 2382 1949 82 40 929 624 819 44 Inv. ex. 2269 1869 8235 918 618 806 45 Inv. ex. 2155 1880 87 37 919 627 782 46 Inv. ex. 23141964 85 38 912 639 813 47 Inv. ex. 2220 1892 85 37 919 610 798

TABLE 1-7 Billet Patenting Quenching Wire break etc. Max. Prior heatingtemp. heating temp. heating temp. [good = no spherical carbide austenitegrain Residual austenite Ex. (° C.) (° C.) (° C.) abnormality] diameter(μm) size (γ#) (vol %) 48 Comp. ex. 1250 930 950 good 0.31 12 9 49 Comp.ex. 1250 930 950 good 0.02 13 7 50 Comp. ex. 1250 930 950 good 0.04 13 851 Comp. ex. 1250 930 950 good 0.12 11 7 52 Comp. ex. 1250 930 950 good0.08 12 21 53 Comp. ex. 1250 930 950 good 0.10 11 2 54 Comp. ex. 1250930 950 good 0.26 13 10 55 Comp. ex. 1250 930 950 good 0.03 10 7 56Comp. ex. 1250 930 950 wire Break — — — 57 Comp. ex. 1250 930 950 wirebreak — — — 58 Comp. ex. 1250 930 950 wire break — — — 59 Comp. ex. 1250930 950 good 0.42 12 7 60 Comp. ex. 1250 930 950 good 0.25 12 8 61 Comp.ex. 1250 930 950 good 0.13 12 6 62 Comp. ex. 1250 930 950 wire break — —— 63 Comp. ex. 1250 930 950 wire break — — — 64 Comp. ex. 1250 930 950good 0.03 11 17 65 Comp. ex. 1250 930 950 good 0.11 11 3 66 Comp. ex.1250 930 950 good 0.33 11 10 67 Comp. ex. 1250 930 950 good 0.03 10 9 68Comp. ex. 1250 930 950 good 0.28 12 11 69 Comp. ex. 1250 930 950 good0.22 13 6 70 Comp. ex. 1250 930 950 good 0.26 12 8 71 Comp. ex. 1250 930950 Decarburized — 12 9 72 Comp. ex. 1250 930 950 Decarburized — 10 1273 Comp. ex. 1100 930 950 good 0.26 10 9 74 Comp. ex. 1100 930 950 good0.28 10 6 Notch Internal Hardness of Tensile bending deg. Nakamura typerotary hardness after nitrided Ex. strength (MPa) 0.2% proof stressYield ratio (%) (deg.) bending (MPa) nitriding (HV) layer (HV) 48 Comp.ex. 2320 1857 80 24 915 617 807 49 Comp. ex. 2002 1806 90 39 812 554 81350 Comp. ex. 2338 1841 79 23 926 612 812 51 Comp. ex. 2236 1954 87 36915 568 728 52 Comp. ex. 2318 1680 75 37 820 612 816 53 Comp. ex. 23951971 82 25 923 637 801 54 Comp. ex. 2363 1909 81 23 920 629 795 55 Comp.ex. 2382 1840 77 39 919 608 731 56 Comp. ex. — — — — — — — 57 Comp. ex.— — — — — — — 58 Comp. ex. — — — — — — — 59 Comp. ex. 2227 1885 85 21929 635 815 60 Comp. ex. 2250 1881 84 17 914 640 816 61 Comp. ex. 23841830 77 17 916 606 817 62 Comp. ex. — — — — — — — 63 Comp. ex. — — — — —— — 64 Comp. ex. 2233 1560 70 21 786 561 783 65 Comp. ex. 2232 1860 8324 911 606 815 66 Comp. ex. 2311 1838 80 21 916 613 783 67 Comp. ex.2276 1977 87 43 789 561 732 68 Comp. ex. 2297 1934 84 24 923 612 809 69Comp. ex. 2243 1838 82 22 914 629 813 70 Comp. ex. 2329 1813 78 21 922596 787 71 Comp. ex. 2359 1898 80 43 774 609 780 72 Comp. ex. 2371 183077 43 781 631 797 73 Comp. ex. 2114 1827 86 25 911 591 793 74 Comp. ex.2251 1822 81 22 910 602 751

TABLE 1-8 Billet Patenting Quenching Wire break etc. Max. Prior heatingtemp. heating temp. heating temp. [good = no spherical carbide austenitegrain Residual austenite Ex. (° C.) (° C.) (° C.) abnormality] diameterμm size (γ#) (vol %) 101 Inv. ex. 1250 — — — 0.03 — — 102 Inv. ex. 1200— — — 0.15 — — 103 Inv. ex. 1250 — — — 0.08 — — 104 Inv. ex. 1250 — — —0.04 — — 105 Inv. ex. 1250 — — — 0.09 — — 106 Inv. ex. 1250 — — — 0.03 —— 107 Inv. ex. 1250 — — — 0.02 — — 108 Inv. ex. 1250 — — — 0.03 — — 109Inv. ex. 1200 — — — 0.14 — — 110 Comp. 1100 — — — 0.21 — — ex. 111 Comp.1100 — — — 0.22 — — ex. Notch Internal Hardness of Tensile bending deg.Nakamura type rotary hardness after nitrided Ex. strength (MPa) 0.2%proof stress Yield ratio (%) (deg.) bending (MPa) nitriding (HV) layer(HV) 101 Inv. ex. — — — — — — — 102 Inv. ex. — — — — — — — 103 Inv. ex.— — — — — — — 104 Inv. ex. — — — — — — — 105 Inv. ex. — — — — — — — 106Inv. ex. — — — — — — — 107 Inv. ex. — — — — — — — 108 Inv. ex. — — — — —— — 109 Inv. ex. — — — — — — — 110 Comp. — — — — — — — ex. 111 Comp. — —— — — — — ex.

INDUSTRIAL APPLICABILITY

The present invention can be utilized for the production of steel wirefor high strength spring use. The high strength spring material can beutilized in many industrial fields starting from the automotiveindustry.

REFERENCE SIGNS LIST

-   1 spherical carbides-   2 punch-   3 test piece-   4 notch-   5 pusher-   6 load use fixture-   P load-   L distance between supports-   θ notch bending angle

1. Pre-drawn steel wire for high strength spring use characterized bycontaining, by mass %, C: 0.67% to less than 0.9%, Si: 2.0 to 3.5%, Mn:0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003 to 0.007%, and Al: 0.0005% to0.003%, having Si and Cr satisfying the following formula:0.3%≦Si−Cr≦1.2%, having a balance of iron and unavoidable impurities,having P and S as impurities comprising P: 0.025% or less and S: 0.025%or less, and, furthermore, having a circle equivalent diameter ofundissolved spherical carbides of less than 0.2 μm.
 2. Pre-drawn steelwire for high strength spring use as set forth in claim 1 characterizedby, further, containing, by mass %, one or more of V: 0.03 to 0.10%, Nb:0.015% or less Mo: 0.05 to 0.30%, W: 0.05 to 0.30% Mg: 0.002% or less,Ca: 0.002% or less, and Zr: 0.003% or less, when containing V satisfying1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and, when containing Mo and W,satisfying 0.05%≦Mo+W≦0.5%.
 3. Drawn heat treated steel wire for highstrength spring use characterized by containing, by mass %, C: 0.67% toless than 0.9%, Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N:0.003 to 0.007%, and Al: 0.0005% to 0.003%, having Si and Cr satisfyingthe following formula:0.3%≦Si−Cr≦1.2%, and having a balance of iron and unavoidableimpurities, having P and S as impurities comprising P: 0.025% or lessand S: 0.025% or less, furthermore, having a metal structure comprisedof at least residual austenite in a volume rate of over 6% to 15%,having prior austenite grain size number of #10 or more, and having acircle equivalent diameter of undissolved spherical carbides of lessthan 0.2 μm.
 4. Drawn heat treated steel wire for high strength springuse as set forth in claim 3 characterized by, further, containing, bymass %, one or more of V: 0.03 to 0.10%, Nb: 0.015% or less Mo: 0.05 to0.30%, W: 0.05 to 0.30% Mg: 0.002% or less, Ca: 0.002% or less, and Zr:0.003% or less, when containing V satisfying 1.4%≦Cr+V≦2.6% and0.70%≦Mn+V≦1.3%, and, when containing Mo and W, satisfying0.05%≦Mo+W≦0.5%.
 5. Drawn heat treated steel wire for high strengthspring use as set forth in claim 3 characterized in that said drawn heattreated steel wire for high strength spring use has a tensile strengthof 2100 to 2400 MPa.
 6. Drawn heat treated steel wire for high strengthspring use as set forth in claim 5 characterized in that said drawn heattreated steel wire for high strength spring use has a yield strength of1600 to 1980 MPa.
 7. Drawn heat treated steel wire for high strengthspring use as set forth in claim 3 characterized in that said drawn heattreated steel wire for high strength spring use has a surface Vicker'shardness of HV750 or more and an internal Vicker's hardness of HV570 ormore aftersoft nitriding of keeping at 500° C. for 1 hour.
 8. A methodof production of pre-drawn steel wire for high strength spring usecharacterized by taking a bloom containing, by mass %, C: 0.67% to lessthan 0.9%, Si: 2.0 to 3.5%, Mn: 0.5 to 1.2%, Cr: 1.3 to 2.5%, N: 0.003to 0.007%, and Al: 0.0005% to 0.003%, having Si and Cr satisfying thefollowing formula: 0.3%≦Si−Cr≦1.2%, having a balance of iron andunavoidable impurities, having P and S as impurities comprising P:0.025% or less and S: 0.025% or less, heating the bloom to 1250° C. ormore, then hot rolling the bloom to produce a billet and heating thebillet to 1200° C. or more, then hot rolling to produce pre-drawn steelwire.
 9. A method of production of pre-drawn steel wire for highstrength spring use as set forth in claim 8 characterized by the bloomfurther, containing, by mass %, one or more of V: 0.03 to 0.10%, Nb:0.015% or less Mo: 0.05 to 0.30%, W: 0.05 to 0.30% Mg: 0.002% or less,Ca: 0.002% or less, and Zr: 0.003% or less, when containing V satisfying1.4%≦Cr+V≦2.6% and 0.70%≦Mn+V≦1.3%, and, when containing Mo and W,satisfying 0.05%≦Mo+W≦0.5%.
 10. A method of production of pre-drawnsteel wire for high strength spring use characterized by further heatingpre-drawn steel wire as set forth in claim 8 to 900° C. or more, thenpatenting it at 600° C. or less.
 11. A method of production of heattreated steel wire for high strength spring use characterized by drawingsaid pre-drawn steel wire which was produced by the method of productionof pre-drawn steel wire as set forth in claim 8, heating it at a heatingrate of 10° C./sec or more up to an A₃ point, holding it at atemperature of the A₃ point or more for 1 minute to 5 minutes, thencooling it at a cooling rate of 50° C./sec or more down to 100° C. orless.
 12. A method of production of heat treated steel wire for highstrength spring use characterized by drawing said pre-drawn steel wirewhich was produced by the method of production of pre-drawn steel wireas set forth in claim 10, heating it at a heating rate of 10° C./sec ormore up to an A₃ point, holding it at a temperature of the A₃ point ormore for 1 minute to 5 minutes, then cooling it at a cooling rate of 50°C./sec or more down to 100° C. or less.
 13. A method of production ofheat treated steel wire for high strength spring use as set forth inclaim 11 characterized by further holding and tempering it at 400 to500° C. for 15 minutes or less.
 14. A method of production of heattreated steel wire for high strength spring use as set forth in claim 12characterized by further holding and tempering it at 400 to 500° C. for15 minutes or less.
 15. Drawn heat treated steel wire for high strengthspring use as set forth in claim 4 characterized in that said drawn heattreated steel wire for high strength spring use has a tensile strengthof 2100 to 2400 MPa.
 16. Drawn heat treated steel wire for high strengthspring use as set forth in claim 4 characterized in that said drawn heattreated steel wire for high strength spring use has a surface Vicker'shardness of HV750 or more and an internal Vicker's hardness of HV570 ormore aftersoft nitriding of keeping at 500° C. for 1 hour.
 17. A methodof production of pre-drawn steel wire for high strength spring usecharacterized by further heating pre-drawn steel wire as set forth inclaims 9 to 900° C. or more, then patenting it at 600° C. or less.
 18. Amethod of production of heat treated steel wire for high strength springuse characterized by drawing said pre-drawn steel wire which wasproduced by the method of production of pre-drawn steel wire as setforth in claim 9, heating it at a heating rate of 10° C./sec or more upto an A₃ point, holding it at a temperature of the A₃ point or more for1 minute to 5 minutes, then cooling it at a cooling rate of 50° C./secor more down to 100° C. or less.